Polymerization-Induced Nano-Order, Ordered Melt, Chain Explosion and Meso-Superstructures of E-Polyethylene

Polymerization-Induced Nano-Order, Ordered Melt, Chain Explosion and Meso-Superstructures of E-Polyethylene | Research Square window.SnipcartSettings = { analytics: { enabled: false } }; (function() { var accessVector = localStorage.getItem('access_vector') || ''; window.dataLayer = window.dataLayer || []; if (accessVector) { window.dataLayer.push({ user: { profile: { profileInfo: { snid: accessVector } } } }); } })(); (function(w,d,s,l,i){w[l]=w[l]||[];w[l].push({'gtm.start':new Date().getTime(),event:'gtm.js'});var f=d.getElementsByTagName(s)[0],j=d.createElement(s),dl=l!='dataLayer'?'&l='+l:'';j.async=true;j.src='https://www.googletagmanager.com/gtm.js?id='+i+dl;f.parentNode.insertBefore(j,f);})(window,document,'script','dataLayer','GTM-K279D39R'); Browse Preprints In Review Journals COVID-19 Preprints AJE Video Bytes Research Tools Research Promotion AJE Professional Editing AJE Rubriq About Preprint Platform In Review Editorial Policies Our Team Advisory Board Help Center Sign In Submit a Preprint Cite Share Download PDF Article Polymerization-Induced Nano-Order, Ordered Melt, Chain Explosion and Meso-Superstructures of E-Polyethylene Aizezi maimaitiming, Kun Cui, Feng Tian This is a preprint; it has not been peer reviewed by a journal. https://doi.org/ 10.21203/rs.3.rs-5814738/v1 This work is licensed under a CC BY 4.0 License Status: Posted Version 1 posted You are reading this latest preprint version Abstract During the past 70 years, it has been well established that melts of crystalline homopolymers adopt disordered random coil conformations and are densely packed in globules above the melting temperature ( T m ). Flow-induced melt orientation promotes crystal nucleation and polymorphism upon cooling. Here, we present polymerization-induced nano-to-mesoscale condensed-ordering as well as order-to-order and order-to-disorder (O-D) melt transformations at T O-D in engineering polyethylene (E-PE). Moreover, we demonstrate fast chain explosion and re-entanglements at T O-D . Liquid-crystalline-like macroscale 2D-ordered nanofibrils and meso-ordered superstructures self-assemble from high-order condensed states and ordered melts via isothermal crystallization, respectively. Self-reinforced optical and shape memory sheets are made by preserving nascent meso-ordering. Eplastomers, which exhibit both elastomeric and soft‒plastic properties, are prepared by nano-fibrillation of non-crosslinked E-PE meso-particles in a polyolefin elastomer matrix and dynamic exchange of 1D-ordered nanofibrillar networks with disordered 3D meso-granules. These findings enable further development of ultrahigh-performance intelligent, photonic and sustainable polymeric materials. Physical sciences/Materials science/Nanoscale materials/Molecular self-assembly Physical sciences/Nanoscience and technology/Nanoscale materials/Structural properties Physical sciences/Chemistry/Polymer chemistry/Mechanical properties Physical sciences/Chemistry/Polymer chemistry/Nanocomposites Physical sciences/Physics/Atomic and molecular physics/Macromolecules and clusters Figures Figure 1 Figure 2 Figure 3 Figure 4 Introduction Compared with the world’s top two commodity plastics of polyethylene (PE) and polypropylene (PP), engineering polyethylene (E-PE) possesses outstanding mechanical, optical, thermal and tribological properties. Advanced structural and functional materials have been made from ultrahigh-molecular-weight polyethylene (UHMWPE) resins. These materials include but are not limited to membrane separators for lithium-ion batteries 1,2 , high-strength fibers/tapes 3–7 , total knee and hip arthroplasty 8,9 and wear resistance modifiers 10 . Potential materials, such as intelligent shape memory fibers 11 , highly thermal conductive nanofibers/films 12,13 , optical transparent high-strength films 13–15 , light actuators 16 , radiative cooling nano-composite film 17 and melt-mechanical recyclable self-reinforced mono-material composites (SMCs) 18,19 , have also been prepared via E-PE. However, owing to the high number of chain entanglements (topological constraints) in the melt, the manufacturing method of UHMWPE fibers and films has been based mainly on solution processing 1–6,11–13,15–17 , which is harmful to the environment and health and, at the same time, is less efficient 7,14, 18–21 . Hierarchical control of the topological constraints and intrinsic morphology of nascent powders (NPs) and crystalline superstructures during polymerization and processing has been vital in the preparation of advanced E-PE materials. For manufacturing high-strength UHMWPE fibers and films, the number of chain entanglements is reduced via continuous solvent extrusion 3,4,13, 15–17 . 1D macro-fibrils of 378–1512 nm, 1D microfibrils of 21–73 nm, and 2D chain-extended crystals of 800 nm were generated via gel spinning and melt drawing 5,6 . These hierarchical structures correlate with the high modulus (65–246 GPa) and high strength (2–7.15 GPa) of UHMWPE fibers 3–7 . The 2D nanofibrillar and microporous superstructures of E-PE films were also constructed via solution gelation and subsequent uniaxial (or biaxial) melt drawing 1,2,13,15 . Owing to the development of (non)metallocene catalysts, “disentangled” NPs have been synthesized, and the solvent-free fabrication of high-strength tapes/films has attracted considerable attention 7,14,18–20 . High-strength/modulus tapes (3.75/157 GPa) and films (0.8/80 GPa) were prepared via solid-state processing below the melting temperatures ( T m ) of the NPs 7,14 . During solution-based and solvent-free melt processing, the morphology of the NPs and processing temperature ( T p ), among many variables, are the most significant factors that strongly influence the performance of E-PE materials 3,7,14,20,22–27 . More importantly, in recent decades, understanding the hierarchical solid-state superstructures and topological constraints of NPs and the low drawability of melt-crystallized E-PEs has been a bottleneck in further development of ultrahigh-performance E-PE materials. Tailoring the ordered melt and crystal structures is of practical importance for preparing membrane separators for lithium-ion batteries 28 , ion conductive films for fuel cell batteries 29,30 , photonic elastomers 31 , and directed self-assembly of sub-10-nm lamellar patterns of block copolymers for microelectronics 32 . High-strength photonic elastomers were prepared by maintaining the multilayered meso-melt of olefin multiblock copolymers (OBCs) below the temperature of order-to-disorder melt transformation ( T O-D ) 31 . By generating shish-kebab crystals via shear-induced orientation, SMCs were formed via injection molding (IM) of melt-blended high-density PE (HDPE) or OBC with a reactor blend of UHMWPE/wax 33 . The E-PE nanofibrillar-toughened PP was made by melt dispersing the intra-crosslinked NPs and compatibilizing them in the PP matrix via a co-rotating four-screw extruder 34 . However, the E-PE NPs cannot be completely dispersed in the matrix above T O-D . The aim of this work is to synergistically control the high-order condensed solid-states and ordered melt, hierarchically dynamically entangled meso-/nanofibrillar networks and multilevel meso-/nanofibrillar superstructures of E-PE. High-strength photonic and intelligent shape memory sheets, re-processable and melt-recyclable dynamic self-reinforced mono-material eplastomers and ultrahigh-impact resistance materials are to be developed under directed processing conditions and crystallization processes. Results Polymerization-induced 2D/3D Self-Assembly and Ordering Figure 1 a and Extended Data Fig. 1 show the 2D/3D self-assembled hierarchical morphologies ranging from macroscale powders and microscale particles to meso-particles and meso-/nanofibrils in the E-PE NPs, which were investigated by scanning electron microscopy (SEM) and ultrasmall-angle X-ray scattering (USAXS). The less entangled grade G1 powders are composed of only non-crosslinked 3D meso-particles with an average size of 886 ± 356 nm. The G2 powders have unconventional heterogeneous meso-/nanofibrillar morphologies. The aggregates of long meso-fibrils with an average width of 606 ± 62 nm are highly intra-crosslinked with long 2D nanofibrils with an average width of 32.4 ± 12.9 nm and length of 2.37 ± 0.83 µm. In addition, meso-particles and separated and aggerated spirals of meso-fibrils are also observed. The G3 powders consisted of homogeneous meso-particles, which were intra-crosslinked with long nanofibrils with an average length of 3.25 ± 1.2 µm. The G4 powders have complicated heterogeneous particle morphologies. The microparticles with average diameters of 11.5 ± 5.7 µm and 22.6 ± 8.1 µm are intra-crosslinked by highly entangled long nano fibrils with an average length of 3.6 ± 1.8 µm. The microparticles contained meso-particles with an average diameter of 596 ± 139 nm. The meso-particles are intra-crosslinked via short nanofibrils with average widths and lengths of 29.7 ± 7.9 nm and 234.2 ± 87.7 nm, respectively. As shown in Extended Data Fig. 2 a, the NPs pressed at temperatures of 30/85/120/128°C have long-term nanostructures with a wide range of statistical domain dimensions between 2.3 nm and 50 nm 22,35 , indicating that chain-extended nanofibrillar crystals may exist in the NPs. As shown in Fig. 1 b and Extended Data Fig. 2 b, the meso-particles and meso-fibrils are in high-order condensed solid-states, exhibiting a high T m of 140–143.1°C and high crystallinity of 71–82.2%. Order-to-Disorder Melt Transformation As shown in Fig. 1 b-c and Extended Data Fig. 3 , the transformation from a high-order condensed solid-state to an ordered melt state occurs after the annealing of the G1-2 powders at an annealing temperature ( T a ) between T m and T O−D of 200°C for 1–2 h. The ordered melt is deduced from both the crystallization temperature ( T c ) of 116.5–119.5°C and the T m of 132–135.4°C. When annealed at T a for 2 h above T O−D , the order melt begins to transform into less-ordered 3D meso-particles (Extended Data Fig. 4 ). When the T a is increased to 300–400°C, the T m of G1 decreases from 134.2°C to 126.4°C, and the T c decreases from 116.2°C to 111.2°C, but the crystallinity ( χ c ) increases from 55.2–74.2%. The low T m of the NPs annealed above T O−D is in accordance with the low T m of the chain-folded crystals of PE according to the Gibbs–Thomson Eq. 3 6 . Compared with less entangled G1, the intrinsic chain entanglements in the NPs strongly promoted the non-isothermal crystallization of G2 (Fig. 1 b and Extended Data Fig. 3 ). Figure 1 c shows the dynamic transition from a high-order condensed solid-state to an order melt state (crystal melting), chain entanglement after melting, chain explosion and re-entanglement upon order-to-disorder melt transition and chain relaxation during shearing. These phenomena are investigated via a torque rheometer, by which the correlation of torque values (melt viscosity) with a wide range of T p values from 130°C to 255°C is quantitatively established. These phenomena are universal for the order melt of all series of NPs, regardless of the molecular weight and intrinsic morphology of the E-PE. The completely molten G3-NPs have the lowest melt viscosity and chain entanglement states at 153°C. When T p increases from 153°C to 172°C, diffusion and coalescence among the molten meso-particles take place, and the chains entangle each other, resulting in a moderate increase in the melt viscosity. These results are completely in contrast to traditional chain explosion and entanglement upon melting 25,37,38 . A continual increase in T p to 190°C induces fast chain explosion and re-entanglement upon order-to-disorder melt transformation. The melt viscosity increased to an un-processible level at 203°C. Above 235°C, the melt viscosity gradually increased to the highest steady value. After shearing for 22–32 min, chain relaxation and disentanglement cause a decrease in the melt viscosity, which is restored to the initial chain-entangled order melt state at 167°C. As shown in Extended Data Fig. 5, chain re-entanglement and relaxation are also proven via melt- and solid-state tensile stress–strain testing, linear melt–rheology and non-isothermal crystallization. The chain re-entanglement and relaxation strongly influence the melt strength, solid-state tensile yielding and neck development, strain hardening property and initial non-isothermal crystallization. The lower yield of the re-entangled G3 sheet is completely distinct from the strong yield of the entangled G3 and HDPE sheets 39 . 2D Self-Assembled Nanofibrillar Superstructures Figure 2 a and Extended Data Fig. 6a present the liquid-crystalline-like macroscale 2D-ordered nanofibrillar superstructures in the compressed G4 sheets. The hierarchical superstructures are self-assembled via compression molding of G4 powders near below the T m . At a compression mold temperature ( T cm ) of 142°C, the transition state from a high-order condensed solid-state to an ordered melt results in the lowest melt viscosity without chain entanglement (Fig. 1 c and Extended Data Fig. 6b-c). Then, macroscale in-plane 2D-ordered nanofibrils with an average width of 28.5 ± 8.5 nm are generated in the G4 sheet via isothermal crystallization at a T c of 128°C for 5–10 h. The USAXS results in Fig. 2 b prove mesoscale ordered crystal superstructures. The quantitative 1D-USAXS curves show long-periodic meso-structures with a wide range of dimensions of 64.9–101.5 nm and their aggregates. The qualitative 2D-USAXS patterns present hierarchical ordered anisotropic superstructures both in the in-plane and through-plane directions, which is also proven by the high T m of G4, but they cannot be measured via 2D-SAXS (Extended Data Fig. 6d-f). Unexpectedly, the long-period dimensions are much greater than the 34.7–50 nm dimensions of the NPs compressed at a T cm of 128°C (Extended Data Fig. 2 a). As shown in Extended Data Fig. 7, the isothermally crystallized G1-3 powders pressed at 140°C also show highly ordered anisotropic meso-morphology both in the in-plane and through-plane directions. Long-period structures with a wide range of domain dimensions of 71.7–88.8 nm, 62.8–80.8 nm and 53.7–64.3 nm are observed for the G1-3 disks, which is in accordance with their high T m values (Extended Data Fig. 8a). In comparison, the microscale 2D-ordered nanofibrils of 22.9 ± 4.1 nm and macroscale randomly 3D-organized nanofibrils self-assembled by pressing at a T cm of 145–220°C followed by isothermal crystallization at a T c of 128°C for 3–10 h. The macroscale randomly 3D-distributed nanofibrils of 23.9 ± 4.7 nm self-organized by quenching from the T cm of 190°C. The long nanofibrils and their aggregates are 3D inter-crosslinked by short nanofibers of 71–101 nm and knots. The tensile properties of unoriented G4 sheets strongly depend on the entangled networks at elevated T cm and microscale 2D-ordered and randomly 3D-organized nanofibrillar superstructures (Extended Data Fig. 6c and Fig. S1 ). The disentangled nanofibrillar superstructure has the lowest ductility (Fig. S1 -2). As shown in Fig. 2 c, the thick G4 bars have the highest impact strength of 202 ± 16.6 KJ/m 2 among all the synthetic polymers. The G5 bars with a high content of nanofibrils exhibit a moderate impact resistance of 101 ± 12 KJ/m 2 , which is higher than the 83.3 ± 1.2 KJ/m 2 of the G3 sheet and the extremely low impact strength of iPP. This result indicates that there is a threshold concentration of hierarchical meso-particles or meso-/nanofibrillar networks that strongly influence the orientation of nanofibrillar aggregates and absorb energy during impact deformation (Extended Data Fig. 6a/8b). For the sheets from the G5 powders with a relatively high concentration of intrinsic nanofibrils, the nanofibrillar networks cannot be easily deformed, the networks may break, and only discontinuous bands of deformation are generated. However, the impact resistance of the G5 sheets is still high enough because of their homogenized nanofibrillar structures. Transition from 2D-Order Meso-fibrils to Disorder 3D Granules Hierarchical order-to-disorder melt transformations during vacuum annealing are shown in Fig. 3 a. The transformation from intra-crosslinked 2D meso-/nanofibrils to disordered 3D granules was observed after the annealing of the NPs at 250°C for 6–20 h under vacuum (Extended Data Fig. 4 a). The reprocessed and isothermal crystallized G2 sheets have T m values of 135.4°C and 130.8°C, respectively, according to the first and second heating curves obtained via DSC. The quenched G2 sheet exhibited interesting ultrahigh tensile ductility, with an average tensile strain at break of 1532 ± 95%, a yield strength of 29.38 ± 0.72 MPa and a low tensile strength at break of 31.73 ± 1.85 MPa. Isothermal crystallization does not significantly influence the tensile properties of G2 (Fig. 3 b). Unexpectedly, the crystallinity of these sheets is as high as 70%. The G2 sheet shows completely distinct tensile deformation behaviors, with strong tensile yielding, long neck development and much lower tensile strain hardening than its inter-crosslinked counterparts. In contrast, the disentangled disordered G2 sheet shows tensile brittleness at a deformation temperature of 120°C (Fig. 3 c). In comparison, the vacuum-annealed low-molecular-weight G3 powders tend to coalesce quickly and undergo chain entanglement, maintaining their order melt (Fig. 3 a). The vacuum-annealed high-molecular-weight G6 powders containing high contents of intrinsic nascent nanofibrils do not compact each other (Extended Data Fig. 4 b). After vacuum annealing, the compressed G3 sheets exhibited high tensile ductility after room-temperature and high-temperature tensile deformation (Fig. 3 b-c and Extended Data Fig. 5f). The G3 sheets show much lower melt drawability, a low melt strength of 0.014 MPa, and a dual-shaped recovery of 48%. After static re-entangling at 220°C for 5 min, the G3 sample has a higher melt strength of 0.055 MPa and a dual-shaped recovery of 72.7% (Fig. 3 d and Extended Data Fig. 5g). As shown in Fig. 3 a, isothermal crystallization at 128°C induced a transition from an ordered melt to a higher-order condensed solid-state in the G3 sheet. In contrast, isothermal crystallization at 116–124°C does not result in a transition from a disordered melt to a condensed ordered solid-state. Meso-Superstructures from Inter-Crosslinked Order Melts Annealing at 142–300°C under pressure suppresses the order-to-disorder melt transition and chain explosion. As shown in Fig. 3 a and Extended Data Fig. 7/8c, the isothermal crystallized and quenched G1-2 sheets, even when pressed at a T cm of 300°C, maintain the ordered melt state. As shown in Fig. 3 b, the isothermal crystallized G2 sheet from 300°C has an average high tensile strength of 75.6 ± 2.4 MPa, similar to the 76.6 ± 3.1 MPa of the G1 sheet. The isothermal crystallized and quenched G2 sheets from 142°C show higher tensile strengths of 63.1 ± 3.2 and 49.5 ± 25 MPa than the 48.8 ± 2.1 and 42.4 ± 1.8 MPa of the G1 sheets (Fig. 3 b, Extended Data Fig. 8d and Table S1 ). The G2 sheets compressed at 142°C also exhibited much greater tensile drawability at the testing temperature of 120°C, with strong strain-hardening properties (Fig. 3 c). The inter-crosslinked G2 sheet has a high melt strength of 0.367 MPa and melt–tensile ductility (Fig. 3 d). These results show that the tensile properties of G2 sheets from 2D meso-fibrils do not depend on the processing temperature or crystallinity due to hierarchical meso-fibrillar and nanofibrillar networks. As shown in Fig. 3 d, the inter-crosslinked G2 sheets show excellent shape recovery of 116% (Extended Data Fig. 5g). Moreover, the isothermal crystallized G2 sheet from 300°C possesses unusual optical transparency and unique selective transmission of visible light. The isothermally crystallized and quenched G2 sheets with thicknesses of 1.1/1.53 µm show high transmittances of 79.8/84.2% and 31/34% for near-infrared and ultraviolet light, respectively, at wavelengths of 1600 and 350 nm (Fig. 3 e). This unique optical transparency was observed only in OBC with a layered mesophase morphology and rod-like lamellar crystals 31 . The qualitative 2D-USAXS patterns in Fig. 3 e-f present hierarchical ordered anisotropic superstructures in both the in-plane and through-plane directions. The quantitative 1D-USAXS curves show that the isothermally crystallized G2 sheet from 300°C has meso-scale superstructures with long periodic dimensions of 49.6–235.6 nm. The quenched G2 sheet from 300°C has a nanoscale superstructure with long periodic dimensions of 36.6–49.5 nm, which is similar to that of the NPs. The isothermally crystallized G1-2 sheets from 142°C have meso-scale superstructures with long periodic dimensions of 88.8–137.7 nm and 44.6–182.6 nm (Fig. 3 f and Extended Data Fig. 7), respectively. The G2 sheets with 2D-oriented nanofibrils and the G4 sheets with 2D-ordered nanofibrils have no strongly oriented orthogonal crystal-stems (Extended Data Fig. 8e and Fig. S3). The orientation parameter of the crystal planes (A [200] /A [110] ) is defined as the peak area ratio of the [200] and [110] reflections of the orthogonal crystals in E-PE. The A [200] /A [110] values of the 2D-oriented G2, 2D-ordered G4 and 3D-random G4 nanofibrils are 0.5, 0.29 and 0.09, respectively, which are much lower than those of high-strength tapes 7 . Eplastomers via Dynamic Nanofibrillar Networks As shown in Fig. 4 a and Extended Data Fig. 9, the eplastomers, composed of hierarchical entangled 5 wt% G1 nanofibrillar networks in the polyolefin elastomer (POE) matrix, possess self-reinforced soft–plastic properties after extrusion and compression molding below T m . The nanofibrillar composite exhibited a high yield strength of 22.9 ± 4.35 MPa, which was much greater than the value of 3.61 ± 0.1 MPa for POE CM at a fixed strain of 50%. The extruded and injection-molded eplastomer (E2 + IM) at a T p of 160°C also behaves as a soft plastic with a high tensile yield strength of 13.93 ± 0.34 MPa due to the shear-induced orientation of the nanofibrils, which is much greater than the 5.35 ± 0.09 MPa of POE IM at a fixed strain of 150%. After compression mold at a T p of 160°C, the epalstomer exhibited elastomeric properties due to the transformation of 2D-ordered nanofibrillar networks to disentangled disordered 3D meso-/nanoparticles. The meso-particle composites have a low tensile strength of 4.23 ± 0.08 MPa at a fixed strain of 50%, a high tensile strength at break of 22.24 ± 0.81 MPa and a tensile strain of 1398 ± 40%. The re-extruded and re-compressed nanofibrillar composites (E3 + CM3) also exhibited soft–plastic properties (Fig. S4). As shown in Fig. 4 b-e and Extended Data Fig. 9, the nanofibrillar crystals have high T m values ranging from 132.3°C to 158°C (Fig. 4 b). The DSC curves show that the non-crosslinked high-order G1 meso-particles can be completely dispersed in the POE matrix (Fig. 4 b). At temperatures of 115°C and 125°C below the T m of the nanofibrils, the compressed nanofibrillar composites have a much higher complex viscosity of 3.22 ~ 4.25 × 10 4 Pa·s than does the 7.65 × 10 3 pure POE at a fixed angular frequency of 0.1 rad/s (Fig. 4 c). The nanofibrillar networks have a 1.97/2.89 times greater diffraction peak intensity for the [110] phase than for the disordered meso-particles, regardless of the compression or injection mold processes (Fig. 4 d). The qualitative 2D-USAXS patterns show that the eplastomers are different from the SMCs of the HDPE/E-PE and OBC/E-PE composites reinforced with shish-kebab crystals 33 . Figure 2 e shows the toluene-etched morphologies of the compressed and injection-molded eplastomers. Homogeneously dispersed nanofibrillar networks are observed. The nanofibrils had an average diameter of 51.1 ± 23.8 nm. Interestingly, the width of these nanofibrils is equal to that of traditional lamellar crystals and the radius of gyration 40 , but the length of the nanofibrils is unexpectedly greater than 10–50 µm. The disordered meso-/microparticles in the POE1-2 matrix have a low melting peak at 127°C (Fig. 4 b), which is in line with the low T m of 126.4°C for high-temperature annealed disordered G1-NPs. The meso-particle composites had slightly greater complex viscosities than did pure POE at low angular frequencies (Fig. 4 c). This result shows that the eplastomers are re-processible and can be melt-mechanically recycled into high-strength soft plastics. The compressed meso-particle composites contain separated and less aggregated meso-granules with average diameters of 640–881 nm and microgranules with average diameters of 1.24–1.95 µm (Fig. 4 e). The abovementioned results prove that the non-crosslinked meso-particles of G1 have potential applications in reinforcing elastomers, rubber and soft plastics and replacing inorganic fillers. In summary, polymerization-induced nano/meso-ordering determines the order of melt below T O−D , order-to-disorder melt transformation at T O−D , chain explosion and re-entanglements at T O−D and meso-order condensed crystalline superstructures from the ordered melt. The order-to-disorder melt transition also significantly depends on the morphology, molecular weight and processing pressure of the NPs. Meso-order condensed superstructures result in unique optical properties. The use of heterogeneous entangled and inter-crosslinked networks results in excellent dual-shaped memory properties. However, homogeneous re-entangled networks lead to low melt ductility. Declarations Data availability The data that support the findings of this study are presented in the paper and/or the Supplementary Information and are available from the corresponding authors on request. Acknowledgments The founding of this research is supported by the National Natural Science Foundation of China (Grant No. U19B6001) and the National Key R&D Plan of China (Grant No. 2020YFC1107003). We thank the beamline BL10U1 station of the Shanghai Synchrotron Radiation Facility for 2D-USAXS measurements. We also thank Yonglong Wu, Xiaoliang Wang and Shunbing Wang for high-temperature solid-state tensile stress–strain testing. Author contributions A.M. developed the original idea, designed and conducted all the experiments, developed the methods, and calculated and analyzed all the results; K.C. provided processing equipment and performed the rheological testing; F.T. performed the 2D-USAXS experiments. A.M. wrote the initial manuscript, revised it, and edited various subsequent versions. Competing interests The authors declare that they have no competing interests. Additional information Supplementary information The online version contains supplementary material available at https://doi.org/X/X. Correspondence and requests for materials should be addressed to Aizezi Maimaitiming. Online Content Any methods, additional references, Nature Portfolio reporting summaries, source data, extended data, supplementary information, acknowledgments, peer review information; details of author contributions and competing interests; and statements of data and code availability are available at https://doi.org/XXX. References Marie, F. L., Raphael, Z. & Vanessa, W. Characterization and performance evaluation of lithium-ion battery separators. Nat. Energy 4, 16–25 (2019). Pankaj, A. & Zhang, J. Z. Battery separators. Chem. Rev. 104, 4419 – 4462 (2004). Wang, J. & Smith, K. J. J. The breaking strength of ultra-high molecular weight polyethylene fibers. Polymer 40, 7261–7274 (1999). Yasuo, O., Hiroki, M. & Takeji, H. 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Methods Materials Commercial nascent G1–6 powders were purchased from Ticona Co., with the following grade names: GUR ® 2024 (G1), GUR ® 4120 (G2), GUR ® 2122 (G4) and GUR ® 4170 (G5). The viscosity-average molecular weights ( M η ) of G1−5 are 5.4E6, 4.2E6, 6.9E5, 4.7E6 and 1.02E7 g/mol (Margolies’ equation). The polyolefin elastomers (POE1−2) were purchased from Dow Chemical Co. The chemical structures of POE1 and iPP are shown in ref. 41, and the grade names of POE2 and iPP are Engage TM 8440 and PPT30S, respectively. The melt flow indices of POE1−2 are 5 and 1 g/10 min (190 ° [email protected] kg), respectively. The antioxidant 1010 was purchased from BASF Co. Characterization and Testing Hierarchical Superstructures : The morphologies of the nascent powders, nanofibrils, meso-particles and fractured surfaces of the sheets were characterized by scanning electron microscopy (SEM, 2 − 10 keV, 5 − 10 mA, HITACHI Regulus 8100). Before characterization, the dumbbell-shaped samples and sheets were cryo-fractured in liquid nitrogen. The surfaces of nascent powders and cryo-fractured sheets were doped with platinum via a HITACHI MCI000. Before the characterization of morphologies, the cryo-fractured surfaces of eplastomers (95POE1 − 2/5G1) emerged in xylene at 80°C for 2–6 h to etch the POE matrix phase. The samples were subsequently dried at 60°C for 6–24 h. Quantitative analysis was performed via Image-Pro Plus 6.0 software. Static transformations from condensed ordering states to order-to-disorder melts of nascent G1–3 powders are tested via differential scanning calorimetry (DSC2500, TA Instrument). Pieces of 4–8 mg of powder were heated from 25°C to an annealing temperature ( T a ) of 140–400°C at a heating rate of 10°C/min, isothermal for 2–5 h, subsequently cooled to 40°C at a cooling rate of 10°C/min, isothermal for 3 min, again heated to an annealing temperature of 160°C, isothermal for 5 min, and cooled to 25–40°C at a cooling rate of 10°C/min. The melting temperature ( T m ) and crystallization temperature ( T c ) are obtained from the heating and cooling curves, respectively. The order-to-disorder melt transformation temperature (T O−D ) is characterized via a sequential annealing and a subsequent non-isothermal crystallization and melting (SA-SC-M) procedure via differential scanning calorimetry (DSC2500, TA Instrument). A 4–8 mg sample of powder or sheet was heated from 25–40°C to an annealing temperature ( T a ) at a heating rate of 10°C/min, isothermal for 1 h, subsequently cooled to 25/40°C at a cooling rate of 10°C/min, isothermal for 3 min, again heated to the annealing temperature of T a + 50°C, and the subsequent crystallization procedure and sequential annealing procedure were repeated under the same conditions. T m and T c are obtained from the heating and cooling curves, respectively. Crystal melting, chain entanglements above T m , chain explosion, re-entanglements and relaxation above T O−D are measured in situ via a torque rheometer (HAAKE PolyLab OS Rheo-Drive 4, Thermo-Scientific Company). The rotation speed is 30–60 rpm. The shearing temperature increases continuously from 130°C to 250°C at intervals of 5°C for 2–4 min. The data acquisition of the torque values is manual from the rheometer software (Extended Data Fig. 10a). The condensed ordering phases of the pressed sheets are analyzed via two-dimensional ultrasmall-angle X-ray scattering experiments (2D-USAXS, λ = 0.124 nm, 10 keV), which are performed at the beamline BL10U1 station of the Shanghai Synchrotron Radiation Facility . The exposure time was 2–20 s. A Dectris Eiger 4 M with a pixel size of 75 × 75 µm was used to collect the signals. The distance between the sample and the detector was 27600 mm, which was calibrated via a bovine tendon standard sample. The scattering data are background corrected according to a general procedure. The crystal stems and lamellar crystals were analyzed via two-dimensional small-/wide-angle X-ray scattering experiments (2D-SAXS/WAXS, λ = 0.124 nm, 10 keV), which were performed at the Shanghai Synchrotron Radiation Facility . The exposure time was 10 s. A Pilatus 2 M detector (Dectris Swiss) with a pixel size of 172 × 172 µm was used to collect the signals. Crystal stems and orientation were also characterized via X-ray powder diffraction (XRD). XRD patterns were obtained with CuKa ( λ = 0.15418 nm) radiation via an X-ray diffractometer (X’Pert Powder, PAN analytical), performing a continuous scan at diffraction angles of 2 θ = 3–50° at 40 mV and 40 mA. The quantitative orientation ratio of the [200] and [110] crystal phases of the orthogonal crystal stems of E-PE is calculated as follows 7 : $$\:\frac{{A}_{\left[200\right]}}{{A}_{\left[110\right]}}=100\:\times\:\:\frac{{A}_{\left[200\right]}}{{A}_{\left[110\right]}}$$ where A [200] and A [110] are the intensities of the [200] and [110] peaks of the orthogonal crystal stems at 2 θ = 21.5°/21.7° and 23.9°/24.1°, respectively, in the XRD profiles. The crystallinity is calculated as follows: $$\:{\chi\:}_{c}=\frac{{\varDelta\:H\times\:100}_{\:}}{{\varDelta\:H}_{PE}^{*}}$$ where ∆ H is the melting enthalpy in the heating DSC curves of the samples. ∆ H PE * is the melting enthalpy of totally crystallized polyethylene with a value of 290 J/g 41 . Solid-state tensile stress–strain tests of dumbbell- and bar-shaped samples cut from the sheets were performed according to the standard methods of ISO527.2-5B/50, ISO20573-TypeCP/50 or ISO527.2-5A/50 (Instron 5966 and Instron 34TM-10). The average values of at least 3–6 tests with high reproducibility were recorded in the tensile strain–stress curves. High-temperature solid–state tensile stress–strain tests of dumbbell-shaped samples cut from the sheets were performed at a testing temperature of 120°C according to the standard methods of ISO527.2-5B/5 (CMT4000, Shenzhen SAS Test Equipment Co. Ltd.). The average values of at least 2–3 tests with high reproducibility were recorded in the tensile stress–strain curves. The melt-state tensile stress–strain properties of the UHMWPE sheets were measured via an ARES-G2 rheometer (TA Instruments). A sheet sample with a size of 35 × 2 × 2 mm was cut from the sheets and compressed and molded at distinct temperatures. The testing temperature was 150°C. The soaking time for every annealing temperature prior to testing was 120 s. The constant linear rate was 2.5–6 mm/min. The dual shape memory effects of the UHMWPE sheets were measured via an ARES-G2 rheometer (TA Instruments). A disentangled G3 sheet and an inter-crosslinked G2 sheet with a size of 35 × 2 × 2 mm were cut from the compression-molded sheets. The samples are melt-drawn at a rate of 2.5 mm/min to a strain of 160% at 150°C, and the temporary shape is fixed under stress after cooling to 40°C. The samples were subsequently recovered into their initial shapes by heating to 150/220°C. The G4 samples were re-entangled at 220°C for 5 min before testing. The shape recovery was calculated according to the standard method 41 . Bi-notched Charpy impact testing for UHMWPE sheets was performed using standard bar samples with a size of 100 × 10 × 4 mm, and two notches with 45° ± 1° and R = 0.25 ± 0.05 mm were machined. At least 6 notched bar samples were tested at an impact energy of 50 J (KD-XJJD-50). Notched Izod impact testing for entangled G3 and iPP bars was performed using standard bar samples with a size of 100 × 10 × 4 mm, and one notch with 45° ± 1° and R = 0.25 ± 0.05 mm was machined. At least 4–5 bar samples were tested at an impact energy of 5.5 J, CESAT9050-5.5. The rheological properties of the UHMWPE and 95POE1/5G1 sheets were measured via an Anton-Paar MCR 302 rheometer. A disk sample with a diameter of 25 mm and a thickness of 1–2 mm is used. The strain is 5%, and the gap is 1–2 mm. The angular frequency is 0.1–100 rad/s, and the testing temperatures are 115–220°C. The soaking time for every annealing temperature prior to testing was 2–10 min. Preparation and Processing of UHMWPE Sheets High-temperature compression: Adequate amounts of nascent G1–5 powders within a mold of 100 × 100 × (1–2) mm were pre-compressed at 0.3 MPa and 128 °C for 400 s, pre-compressed at 2.5 MPa and 128 °C for 200 s, compressed at 10 MPa and 128 °C for 900 s, and then increased to 142–145 °C for 400 s under a pressure of 0.3 MPa. The preformed samples were compressed at 2.5 MPa and 142–145 °C for 200 s, the temperature was increased to 300 °C for 30 min under a pressure of 10 MPa, and the sample was compressed at 10 MPa and 300 °C for 2 h. Then, the preformed sheets in the molds were transformed into upper layers of a molding presser at a temperature of 124–128 °C, isothermally crystallized for 3–6 h under a pressure of 10 MPa, and slowly cooled to room temperature in the mold (Extended Data Fig. 10b). Preparation of sheets with 2D-ordered and 3D randomly constructed nanofibrils: Adequate amounts of nascent G1–5 powders in a mold with a size of 100 × 100 × (1–2) mm or 240 × 240 × (2–40) mm were pre-compressed at 0.3 MPa and 128 °C for 200–600 s, pre-compressed at 2.5 MPa and 128 °C for 400–900 s, and compressed at 10 MPa and 128 °C for 300–900 s. Then, the temperature was increased to 142–145 °C or 145–220 °C for 400 s–1 h under a pressure of 0.3 MPa. The preformed samples were compressed at 2.5 MPa and 142–145 °C or 145–220 °C for 200–900 s, and compressed at 10 MPa and 142–145 °C or 145–220 °C for 120 s–2 h. Then, the temperature was decreased to 128 °C for 30 min under a pressure of 10 MPa, or the preformed sheets in molds were transformed into the upper layers of a molding presser at a temperature of 128 °C. The sheets were subsequently isothermally crystallized for 30 min–23 h under a pressure of 10 MPa. For microscale 2D ordering at 145–150 °C, the pre-formed sheets in molds were quenched in the upper layers at 20–60 °C. For macroscale random 3D construction of nanofibrils, the compression molding temperatures were set to 145–220 °C, and the preformed sheets in the molds were quenched in the upper layers at 20–80 °C under a pressure of 10 MPa (Extended Data Fig. 10b). Dynamic Nano-fibrillation of Nascent G1 Powders in POE Matrix Premixing and nano-fibrillation of G1 in POE1–2: A batch internal mixer (HAAKE PolyLab OS Rheo-Drive 4, Thermo-Scientific Company) is used to premix or melt blend G1 with POE1−2. The rotation speed was 60/120 rpm, the mixing temperature was 80/115 °C, and the mixing time was 6 mm (Extended Data Fig. 10a). Compression molding: Pellets or blends are pre-compressed in the mold at 0.3 MPa and 110/160 °C for 300 s, pre-compressed at 0.3 MPa and 110/160 °C for 100 s, pre-compressed at 2.5 MPa and 110/160 °C for 200 s, and compressed at 10 MPa and 110/160 °C for 300 s. The preformed samples are quenched in the upper layers of a molding presser at 25–50 °C for 5 min (Extended Data Fig. 10b). Extruded nano-fibrillation of G1 in POE1: A laboratory-scale corotating twin-screw extruder (HAAKE MiniLab Rheomex CTW5, Thermo-Scientific Company) is used to compound solid-state premixed POE1/G1 under cycling conditions for 5–15 min. The weight composition ratio of G1 to POE is 95:5 wt%. The temperatures of the distinct zones are 115/135 °C. The rotation speed of the motor is 80–300 rpm (Extended Data Fig. 10c). Shear-induced orientation of G1 meso-particles into nanofibrils in POE1–2: A laboratory-scale HAAKE MiniJet Pro combined with a corotating twin-screw extruder (HAAKE MiniLab Rheomex, Thermo-Scientific Company) is used to melt compounding and injection-molding of the pressed POE1-2/G1 nanocomposites. The weight composition ratios of G1 to POE1-2 are (100/99/97.5/95): (0/1/2.5/5) wt%. The extrusion/injection/molding temperatures are 160/190 °C and 40 °C. The rotation speed of the extruder was 80 rpm. References 41. Aizezi, M. et al. High-strength triple shape memory elastomers from radiation-vulcanized polyolefin elastomer/polypropylene blends. ACS Appl. Polym. Mater. 1 , 1735–1748 (2019). Additional Declarations There is NO Competing Interest. Supplementary Files SIAzizNNSM0112.docx Polymerization-Induced Nano-Order, Ordered Melt, Chain Explosion and Meso-Superstructures of E-Polyethylene ExtendedDataFigs.docx Cite Share Download PDF Status: Posted Version 1 posted You are reading this latest preprint version Research Square lets you share your work early, gain feedback from the community, and start making changes to your manuscript prior to peer review in a journal. As a division of Research Square Company, we’re committed to making research communication faster, fairer, and more useful. We do this by developing innovative software and high quality services for the global research community. Our growing team is made up of researchers and industry professionals working together to solve the most critical problems facing scientific publishing. Also discoverable on Platform About Our Team In Review Editorial Policies Advisory Board Help Center Resources Author Services Accessibility API Access RSS feed Manage Cookie Preferences © Research Square 2026 | ISSN 2693-5015 (online) Privacy Policy Terms of Service Do Not Sell My Personal Information {"props":{"pageProps":{"initialData":{"identity":"rs-5814738","acceptedTermsAndConditions":true,"allowDirectSubmit":true,"archivedVersions":[],"articleType":"Article","associatedPublications":[],"authors":[{"id":401961030,"identity":"ed5018dd-e143-4547-a265-b7777c569efd","order_by":0,"name":"Aizezi maimaitiming","email":"data:image/png;base64,iVBORw0KGgoAAAANSUhEUgAAAZAAAAAyAQMAAABI0h/eAAAABlBMVEX///8AAABVwtN+AAAACXBIWXMAAA7EAAAOxAGVKw4bAAAAtElEQVRIiWNgGAWjYFACxgYGhgoGHlK1nCFNC0hXGymq+fsPN3+6Oa9Ohp/9AOOHHwx2eQS1SBw42GCcu+0wj2RPArNkD0NyMWFrDjY2JOduO8BjcIOBQZqB4UBiAyEd8ocZGw7nzqkDaWH+TZQWg2OMjc25DcwgLWzE2WJ4hrGZOecYyC+JbZY9BsmEtcidP/74c05NnT0/++HDN35U2BHWggRAcWpAgvpRMApGwSgYBbgBAL8MOB/no1q8AAAAAElFTkSuQmCC","orcid":"https://orcid.org/0000-0002-6595-0176","institution":"Shanghai Institute of Organic Chemistry, Chinese Academy of Sciences","correspondingAuthor":true,"prefix":"","firstName":"Aizezi","middleName":"","lastName":"maimaitiming","suffix":""},{"id":401961031,"identity":"7e02d93a-55d8-4462-8b5a-588c28b108d6","order_by":1,"name":"Kun Cui","email":"","orcid":"","institution":"Shanghai Institute of Organic Chemistry, Chinese Academy of Sciences","correspondingAuthor":false,"prefix":"","firstName":"Kun","middleName":"","lastName":"Cui","suffix":""},{"id":401961032,"identity":"f1c040fb-8d8c-4711-b3c6-9334732a8e05","order_by":2,"name":"Feng Tian","email":"","orcid":"","institution":"Shanghai Synchrotron Radiation Facility, Shanghai Advanced Research Institute, Chinese Academy of Sciences","correspondingAuthor":false,"prefix":"","firstName":"Feng","middleName":"","lastName":"Tian","suffix":""}],"badges":[],"createdAt":"2025-01-12 16:45:24","currentVersionCode":1,"declarations":"","doi":"10.21203/rs.3.rs-5814738/v1","doiUrl":"https://doi.org/10.21203/rs.3.rs-5814738/v1","draftVersion":[],"editorialEvents":[],"editorialNote":"","failedWorkflow":false,"files":[{"id":73829128,"identity":"20dc8cd3-7040-4dbc-a057-7b5c2278c62d","added_by":"auto","created_at":"2025-01-15 06:02:17","extension":"png","order_by":1,"title":"Figure 1","display":"","copyAsset":false,"role":"figure","size":391763,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003ePolymerization-induced nano/meso-ordering, static and dynamic transformations from high-order condensed solid-state to order-to-disorder melts, and chain explosion and re-entanglements at \u003c/strong\u003e\u003cem\u003e\u003cstrong\u003eT\u003c/strong\u003e\u003c/em\u003e\u003csub\u003e\u003cstrong\u003eO-D\u003c/strong\u003e\u003c/sub\u003e\u003cstrong\u003e of engineering polyethylene\u003c/strong\u003e. \u003cstrong\u003ea\u003c/strong\u003e: Hierarchical non-crosslinked G1 meso-particles of 886 ± 356 nm; G2 meso-fibrils of 606 ± 62 nm are intra-crosslinked with long nanofibrils of 32.4 ± 12.9 nm and 2.37 ± 0.83 μm; G3 meso-particles are intra-crosslinked with long nanofibrils of 3.25 ± 1.2 μm; \u003cstrong\u003eb-c\u003c/strong\u003e: Static transformations from high-order condensed solid-states to order-to-disorder melts of G1-2 powders by annealing at distinct temperatures of 140–420 °C for 1/2 h under atmospheric pressure (1/2 h: melting and crystallinity versus annealing temperatures (\u003cem\u003eT\u003c/em\u003e\u003csub\u003ea\u003c/sub\u003e) during independent/sequential annealing, subsequent crystallization and melting procedures at a standard heating and cooling rate of 10 °C/min); \u003cstrong\u003ed\u003c/strong\u003e: Crystal melting, chain entanglement, fast chain explosion and re-entanglements and slow chain relaxation during dynamic shearing at distinct processing temperatures.\u003c/p\u003e","description":"","filename":"1.png","url":"https://assets-eu.researchsquare.com/files/rs-5814738/v1/e989ba41a17603b0bb07ead5.png"},{"id":73826002,"identity":"56a7dfef-8720-479b-ac13-069bfd65a6d9","added_by":"auto","created_at":"2025-01-15 05:29:39","extension":"png","order_by":2,"title":"Figure 2","display":"","copyAsset":false,"role":"figure","size":440551,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003e2D self-assembled nanofibrillar superstructures, characterization, intrinsic nanofibrillar networks and impact strength. a\u003c/strong\u003e: Macroscale and microscale 2D-ordered (2D-Od) nanofibrils obtained by melting nascent G4 powders at 142 °C and 145–210 °C followed by isothermal crystallization at 128 °C for 3–10 h, and macroscale randomly 3D-constructed nanofibrils by pressing G4-NPs at 190 °C followed by quenched crystallization below 60 °C; \u003cstrong\u003eb\u003c/strong\u003e: Characterization of macroscale 2D-ordered nano-to-mesoscale structures via ultrasmall-angle X-ray scattering (USAXS) (G3-NP-128 °C: nanostructure of the G3 disk compressed at 128 °C for comparison); \u003cstrong\u003ec\u003c/strong\u003e: Bi-notched Charpy and Izod impact strength of random G3-5 sheets with a thickness of 4 mm (iPP: isotactic polypropylene).\u003c/p\u003e","description":"","filename":"2.png","url":"https://assets-eu.researchsquare.com/files/rs-5814738/v1/d07f522738955a4f27653657.png"},{"id":73826023,"identity":"469de470-af50-4407-8697-22a5d2fd4d76","added_by":"auto","created_at":"2025-01-15 05:29:40","extension":"png","order_by":3,"title":"Figure 3","display":"","copyAsset":false,"role":"figure","size":287036,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eHierarchical transformation from intra-crosslinked high-order condensed solid-states to disentangled disorder melts and inter-crosslinked order melts, intelligent shape memory effects, and optical properties of meso-order superstructures. a\u003c/strong\u003e: Static transformation from an intra-crosslinked high-order condensed state to disentangle disordered melt of G2-NPs and an entangled-order melt of G3-NPs via vacuum annealing at 250 °C for 6–20 h (Extended Data Fig. 4) and pressing at 150 °C; static transformation from an intra-crosslinked high-order condensed state to an inter-crosslinked and heterogeneously entangled-order melt of G2-NPs by pressing at 300 °C for 2 h and subsequent crystallization; transition from an ordered melt to condensed ordered solid-states of compressed and vacuum-annealed G2-3 during isothermal crystallization (IC: isothermally crystallized at 128 °C for 3 h from an ordered melt and 116–124 °C for 6 h from a disordered melt; Q: quenched in the mold below 60 °C; solid and dashed lines: first and second heating curves in DSC); \u003cstrong\u003eb-c\u003c/strong\u003e: Room- and high-temperature solid-state tensile stress–strain properties of disentangled disordered G2 sheets, inter-crosslinked order G2 sheets and entangled order G3 sheets; \u003cstrong\u003ed\u003c/strong\u003e: Melt-state tensile stress–strain properties and dual-shaped memory effects; \u003cstrong\u003ee-f\u003c/strong\u003e: Selective visible-light transmittance properties and characterization of condensed ordered solid-states of G2 sheets with thicknesses of 1.1/1.53 μm via USAXS.\u003c/p\u003e","description":"","filename":"3.png","url":"https://assets-eu.researchsquare.com/files/rs-5814738/v1/32133fcd9edf28614a60f2ad.png"},{"id":73826014,"identity":"fcb38b59-6a22-4232-a17b-b5ab13020f1d","added_by":"auto","created_at":"2025-01-15 05:29:39","extension":"png","order_by":4,"title":"Figure 4","display":"","copyAsset":false,"role":"figure","size":357373,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eSelf-reinforced mono-material eplastomers via dynamic nanofibrillar networks. a\u003c/strong\u003e: Tensile stress–strain properties of eplastomers; soft plastic properties via shear-induced orientation of G1 nanofibrillar networks in the polyolefin elastomer (POE) matrix during injection molding (IM) at 160 °C above \u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e and compression molding (CM) at 110 °C below \u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e; elastic properties by melting-induced transformation of the ordered nanofibrils into disordered meso-particles during CM at 160 °C above \u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e; \u003cstrong\u003eb\u003c/strong\u003e: Completely transformation of 5 wt% high-order G1 meso-particles into 1D-ordered nanofibrillar networks in POE via extrusion (E1) at 115/135 °C below \u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e; \u003cstrong\u003ec\u003c/strong\u003e: Rheological properties of dynamic eplastomers containing entangled nanofibrillar networks at 115–135 °C below \u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e and disordered meso-particles at 145–160 °C above \u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e; \u003cstrong\u003ed-e\u003c/strong\u003e: Characterization of 2D-ordered nanofibrils of 51.1 ± 23.8 nm, disordered meso-particles of 640–881nm and microparticles of 1.24–1.95 μm via XRD, USAXS and SEM.\u003c/p\u003e","description":"","filename":"4.png","url":"https://assets-eu.researchsquare.com/files/rs-5814738/v1/557988d31641da41abf55499.png"},{"id":73923869,"identity":"404dbb6d-4904-4973-b6e7-bab09d5e1a71","added_by":"auto","created_at":"2025-01-16 04:18:55","extension":"pdf","order_by":0,"title":"","display":"","copyAsset":false,"role":"manuscript-pdf","size":2622611,"visible":true,"origin":"","legend":"","description":"","filename":"manuscript.pdf","url":"https://assets-eu.researchsquare.com/files/rs-5814738/v1/9b67e9d0-596d-42b1-811d-c9ad43c6db63.pdf"},{"id":73826000,"identity":"b731fd37-145f-49c0-a009-b1f363da2822","added_by":"auto","created_at":"2025-01-15 05:29:39","extension":"docx","order_by":1,"title":"","display":"","copyAsset":false,"role":"supplement","size":1658021,"visible":true,"origin":"","legend":"Polymerization-Induced Nano-Order, Ordered Melt, Chain Explosion and Meso-Superstructures of E-Polyethylene","description":"","filename":"SIAzizNNSM0112.docx","url":"https://assets-eu.researchsquare.com/files/rs-5814738/v1/8f0de88bc5dfb5d54ce00948.docx"},{"id":73826004,"identity":"4fa20f50-964c-4c33-8015-0b82148937fb","added_by":"auto","created_at":"2025-01-15 05:29:39","extension":"docx","order_by":2,"title":"","display":"","copyAsset":false,"role":"supplement","size":17117565,"visible":true,"origin":"","legend":"","description":"","filename":"ExtendedDataFigs.docx","url":"https://assets-eu.researchsquare.com/files/rs-5814738/v1/f63fb4af80c0ee0a422c1bce.docx"}],"financialInterests":"There is \u003cb\u003eNO\u003c/b\u003e Competing Interest.","formattedTitle":"Polymerization-Induced Nano-Order, Ordered Melt, Chain Explosion and Meso-Superstructures of E-Polyethylene","fulltext":[{"header":"Introduction","content":"\u003cp\u003eCompared with the world\u0026rsquo;s top two commodity plastics of polyethylene (PE) and polypropylene (PP), engineering polyethylene (E-PE) possesses outstanding mechanical, optical, thermal and tribological properties. Advanced structural and functional materials have been made from ultrahigh-molecular-weight polyethylene (UHMWPE) resins. These materials include but are not limited to membrane separators for lithium-ion batteries\u003csup\u003e1,2\u003c/sup\u003e, high-strength fibers/tapes\u003csup\u003e3\u0026ndash;7\u003c/sup\u003e, total knee and hip arthroplasty\u003csup\u003e8,9\u003c/sup\u003e and wear resistance modifiers\u003csup\u003e10\u003c/sup\u003e. Potential materials, such as intelligent shape memory fibers\u003csup\u003e11\u003c/sup\u003e, highly thermal conductive nanofibers/films\u003csup\u003e12,13\u003c/sup\u003e, optical transparent high-strength films\u003csup\u003e13\u0026ndash;15\u003c/sup\u003e, light actuators\u003csup\u003e16\u003c/sup\u003e, radiative cooling nano-composite film\u003csup\u003e17\u003c/sup\u003e and melt-mechanical recyclable self-reinforced mono-material composites (SMCs)\u003csup\u003e18,19\u003c/sup\u003e, have also been prepared via E-PE. However, owing to the high number of chain entanglements (topological constraints) in the melt, the manufacturing method of UHMWPE fibers and films has been based mainly on solution processing\u003csup\u003e1\u0026ndash;6,11\u0026ndash;13,15\u0026ndash;17\u003c/sup\u003e, which is harmful to the environment and health and, at the same time, is less efficient\u003csup\u003e7,14, 18\u0026ndash;21\u003c/sup\u003e.\u003c/p\u003e\n\u003cp\u003eHierarchical control of the topological constraints and intrinsic morphology of nascent powders (NPs) and crystalline superstructures during polymerization and processing has been vital in the preparation of advanced E-PE materials. For manufacturing high-strength UHMWPE fibers and films, the number of chain entanglements is reduced via continuous solvent extrusion\u003csup\u003e3,4,13, 15\u0026ndash;17\u003c/sup\u003e. 1D macro-fibrils of 378\u0026ndash;1512 nm, 1D microfibrils of 21\u0026ndash;73 nm, and 2D chain-extended crystals of 800 nm were generated via gel spinning and melt drawing\u003csup\u003e5,6\u003c/sup\u003e. These hierarchical structures correlate with the high modulus (65\u0026ndash;246 GPa)\u0026nbsp;and high strength (2\u0026ndash;7.15 GPa) of UHMWPE fibers\u003csup\u003e3\u0026ndash;7\u003c/sup\u003e. The 2D nanofibrillar and microporous superstructures of E-PE films were also constructed via solution gelation and subsequent uniaxial (or biaxial) melt drawing\u003csup\u003e1,2,13,15\u003c/sup\u003e. Owing to the development of (non)metallocene catalysts, \u0026ldquo;disentangled\u0026rdquo; NPs have been synthesized, and the solvent-free fabrication of high-strength tapes/films has attracted considerable attention\u003csup\u003e7,14,18\u0026ndash;20\u003c/sup\u003e. High-strength/modulus tapes (3.75/157 GPa) and films (0.8/80 GPa) were prepared via solid-state processing below the melting temperatures (\u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e) of the NPs\u003csup\u003e7,14\u003c/sup\u003e. During solution-based and solvent-free melt processing, the morphology of the NPs and processing temperature (\u003cem\u003eT\u003c/em\u003e\u003csub\u003ep\u003c/sub\u003e), among many variables, are the most significant factors that strongly influence the performance of E-PE materials\u003csup\u003e3,7,14,20,22\u0026ndash;27\u003c/sup\u003e. More importantly, in recent decades, understanding the hierarchical solid-state superstructures and topological constraints of NPs and the low drawability of melt-crystallized E-PEs has been a bottleneck in further development of ultrahigh-performance E-PE materials.\u0026nbsp;\u003c/p\u003e\n\u003cp\u003eTailoring the ordered melt and crystal structures is of practical importance for preparing membrane separators for lithium-ion batteries\u003csup\u003e28\u003c/sup\u003e, ion conductive films for fuel cell batteries\u003csup\u003e29,30\u003c/sup\u003e, photonic elastomers\u003csup\u003e31\u003c/sup\u003e, and directed self-assembly of sub-10-nm lamellar patterns of block copolymers for microelectronics\u003csup\u003e32\u003c/sup\u003e. High-strength photonic elastomers were prepared by maintaining the multilayered meso-melt of olefin multiblock copolymers (OBCs) below the temperature of order-to-disorder melt transformation (\u003cem\u003eT\u003c/em\u003e\u003csub\u003eO-D\u003c/sub\u003e)\u003csup\u003e31\u003c/sup\u003e. By generating shish-kebab crystals via shear-induced orientation, SMCs were formed via injection molding (IM) of melt-blended high-density PE (HDPE) or OBC with a reactor blend of UHMWPE/wax\u003csup\u003e33\u003c/sup\u003e. The E-PE nanofibrillar-toughened PP was made by melt dispersing the intra-crosslinked NPs and compatibilizing them in the PP matrix via a co-rotating four-screw extruder\u003csup\u003e34\u003c/sup\u003e. However, the E-PE NPs cannot be completely dispersed in the matrix above \u003cem\u003eT\u003c/em\u003e\u003csub\u003eO-D\u003c/sub\u003e.\u0026nbsp;\u003c/p\u003e\n\u003cp\u003eThe aim of this work is to synergistically control the high-order condensed solid-states and ordered melt, hierarchically dynamically entangled meso-/nanofibrillar networks and multilevel meso-/nanofibrillar superstructures of E-PE. High-strength photonic and intelligent shape memory sheets, re-processable and melt-recyclable dynamic self-reinforced mono-material eplastomers and ultrahigh-impact resistance materials are to be developed under directed processing conditions and crystallization processes.\u0026nbsp;\u003c/p\u003e"},{"header":"Results","content":"\n\u003ch3\u003ePolymerization-induced 2D/3D Self-Assembly and Ordering\u003c/h3\u003e\n\u003cp\u003eFigure\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ea and Extended Data Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003e show the 2D/3D self-assembled hierarchical morphologies ranging from macroscale powders and microscale particles to meso-particles and meso-/nanofibrils in the E-PE NPs, which were investigated by scanning electron microscopy (SEM) and ultrasmall-angle X-ray scattering (USAXS). The less entangled grade G1 powders are composed of only non-crosslinked 3D meso-particles with an average size of 886\u0026thinsp;\u0026plusmn;\u0026thinsp;356 nm. The G2 powders have unconventional heterogeneous meso-/nanofibrillar morphologies. The aggregates of long meso-fibrils with an average width of 606\u0026thinsp;\u0026plusmn;\u0026thinsp;62 nm are highly intra-crosslinked with long 2D nanofibrils with an average width of 32.4\u0026thinsp;\u0026plusmn;\u0026thinsp;12.9 nm and length of 2.37\u0026thinsp;\u0026plusmn;\u0026thinsp;0.83 \u0026micro;m. In addition, meso-particles and separated and aggerated spirals of meso-fibrils are also observed. The G3 powders consisted of homogeneous meso-particles, which were intra-crosslinked with long nanofibrils with an average length of 3.25\u0026thinsp;\u0026plusmn;\u0026thinsp;1.2 \u0026micro;m. The G4 powders have complicated heterogeneous particle morphologies. The microparticles with average diameters of 11.5\u0026thinsp;\u0026plusmn;\u0026thinsp;5.7 \u0026micro;m and 22.6\u0026thinsp;\u0026plusmn;\u0026thinsp;8.1 \u0026micro;m are intra-crosslinked by highly entangled long nano fibrils with an average length of 3.6\u0026thinsp;\u0026plusmn;\u0026thinsp;1.8 \u0026micro;m. The microparticles contained meso-particles with an average diameter of 596\u0026thinsp;\u0026plusmn;\u0026thinsp;139 nm. The meso-particles are intra-crosslinked via short nanofibrils with average widths and lengths of 29.7\u0026thinsp;\u0026plusmn;\u0026thinsp;7.9 nm and 234.2\u0026thinsp;\u0026plusmn;\u0026thinsp;87.7 nm, respectively.\u003c/p\u003e\u003cp\u003eAs shown in Extended Data Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ea, the NPs pressed at temperatures of 30/85/120/128\u0026deg;C have long-term nanostructures with a wide range of statistical domain dimensions between 2.3 nm and 50 nm\u003csup\u003e22,35\u003c/sup\u003e, indicating that chain-extended nanofibrillar crystals may exist in the NPs. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003eb and Extended Data Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003eb, the meso-particles and meso-fibrils are in high-order condensed solid-states, exhibiting a high \u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e of 140\u0026ndash;143.1\u0026deg;C and high crystallinity of 71\u0026ndash;82.2%.\u003c/p\u003e\n\u003ch3\u003eOrder-to-Disorder Melt Transformation\u003c/h3\u003e\n\u003cp\u003eAs shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003eb-c and Extended Data Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003e, the transformation from a high-order condensed solid-state to an ordered melt state occurs after the annealing of the G1-2 powders at an annealing temperature (\u003cem\u003eT\u003c/em\u003e\u003csub\u003ea\u003c/sub\u003e) between \u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e and \u003cem\u003eT\u003c/em\u003e\u003csub\u003eO\u0026minus;D\u003c/sub\u003e of 200\u0026deg;C for 1\u0026ndash;2 h. The ordered melt is deduced from both the crystallization temperature (\u003cem\u003eT\u003c/em\u003e\u003csub\u003ec\u003c/sub\u003e) of 116.5\u0026ndash;119.5\u0026deg;C and the \u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e of 132\u0026ndash;135.4\u0026deg;C. When annealed at \u003cem\u003eT\u003c/em\u003e\u003csub\u003ea\u003c/sub\u003e for 2 h above \u003cem\u003eT\u003c/em\u003e\u003csub\u003eO\u0026minus;D\u003c/sub\u003e, the order melt begins to transform into less-ordered 3D meso-particles (Extended Data Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003e). When the \u003cem\u003eT\u003c/em\u003e\u003csub\u003ea\u003c/sub\u003e is increased to 300\u0026ndash;400\u0026deg;C, the \u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e of G1 decreases from 134.2\u0026deg;C to 126.4\u0026deg;C, and the \u003cem\u003eT\u003c/em\u003e\u003csub\u003ec\u003c/sub\u003e decreases from 116.2\u0026deg;C to 111.2\u0026deg;C, but the crystallinity (\u003cem\u003eχ\u003c/em\u003e\u003csub\u003ec\u003c/sub\u003e) increases from 55.2\u0026ndash;74.2%. The low \u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e of the NPs annealed above \u003cem\u003eT\u003c/em\u003e\u003csub\u003eO\u0026minus;D\u003c/sub\u003e is in accordance with the low \u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e of the chain-folded crystals of PE according to the Gibbs\u0026ndash;Thomson Eq.\u0026nbsp;3\u003csup\u003e6\u003c/sup\u003e. Compared with less entangled G1, the intrinsic chain entanglements in the NPs strongly promoted the non-isothermal crystallization of G2 (Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003eb and Extended Data Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003e).\u003c/p\u003e\u003cp\u003eFigure\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ec shows the dynamic transition from a high-order condensed solid-state to an order melt state (crystal melting), chain entanglement after melting, chain explosion and re-entanglement upon order-to-disorder melt transition and chain relaxation during shearing. These phenomena are investigated via a torque rheometer, by which the correlation of torque values (melt viscosity) with a wide range of \u003cem\u003eT\u003c/em\u003e\u003csub\u003ep\u003c/sub\u003e values from 130\u0026deg;C to 255\u0026deg;C is quantitatively established. These phenomena are universal for the order melt of all series of NPs, regardless of the molecular weight and intrinsic morphology of the E-PE. The completely molten G3-NPs have the lowest melt viscosity and chain entanglement states at 153\u0026deg;C. When \u003cem\u003eT\u003c/em\u003e\u003csub\u003ep\u003c/sub\u003e increases from 153\u0026deg;C to 172\u0026deg;C, diffusion and coalescence among the molten meso-particles take place, and the chains entangle each other, resulting in a moderate increase in the melt viscosity. These results are completely in contrast to traditional chain explosion and entanglement upon melting\u003csup\u003e25,37,38\u003c/sup\u003e. A continual increase in \u003cem\u003eT\u003c/em\u003e\u003csub\u003ep\u003c/sub\u003e to 190\u0026deg;C induces fast chain explosion and re-entanglement upon order-to-disorder melt transformation. The melt viscosity increased to an un-processible level at 203\u0026deg;C. Above 235\u0026deg;C, the melt viscosity gradually increased to the highest steady value. After shearing for 22\u0026ndash;32 min, chain relaxation and disentanglement cause a decrease in the melt viscosity, which is restored to the initial chain-entangled order melt state at 167\u0026deg;C.\u003c/p\u003e \u003cp\u003eAs shown in Extended Data Fig.\u0026nbsp;5, chain re-entanglement and relaxation are also proven via melt- and solid-state tensile stress\u0026ndash;strain testing, linear melt\u0026ndash;rheology and non-isothermal crystallization. The chain re-entanglement and relaxation strongly influence the melt strength, solid-state tensile yielding and neck development, strain hardening property and initial non-isothermal crystallization. The lower yield of the re-entangled G3 sheet is completely distinct from the strong yield of the entangled G3 and HDPE sheets\u003csup\u003e39\u003c/sup\u003e.\u003c/p\u003e \n\u003ch3\u003e2D Self-Assembled Nanofibrillar Superstructures\u003c/h3\u003e\n\u003cdiv id=\"Sec3\" class=\"Section2\"\u003e \u003cp\u003eFigure\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ea and Extended Data Fig.\u0026nbsp;6a present the liquid-crystalline-like macroscale 2D-ordered nanofibrillar superstructures in the compressed G4 sheets. The hierarchical superstructures are self-assembled via compression molding of G4 powders near below the \u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e. At a compression mold temperature (\u003cem\u003eT\u003c/em\u003e\u003csub\u003ecm\u003c/sub\u003e) of 142\u0026deg;C, the transition state from a high-order condensed solid-state to an ordered melt results in the lowest melt viscosity without chain entanglement (Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ec and Extended Data Fig.\u0026nbsp;6b-c). Then, macroscale in-plane 2D-ordered nanofibrils with an average width of 28.5\u0026thinsp;\u0026plusmn;\u0026thinsp;8.5 nm are generated in the G4 sheet via isothermal crystallization at a \u003cem\u003eT\u003c/em\u003e\u003csub\u003ec\u003c/sub\u003e of 128\u0026deg;C for 5\u0026ndash;10 h.\u003c/p\u003e \u003cp\u003eThe USAXS results in Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003eb prove mesoscale ordered crystal superstructures. The quantitative 1D-USAXS curves show long-periodic meso-structures with a wide range of dimensions of 64.9\u0026ndash;101.5 nm and their aggregates. The qualitative 2D-USAXS patterns present hierarchical ordered anisotropic superstructures both in the in-plane and through-plane directions, which is also proven by the high \u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e of G4, but they cannot be measured via 2D-SAXS (Extended Data Fig.\u0026nbsp;6d-f). Unexpectedly, the long-period dimensions are much greater than the 34.7\u0026ndash;50 nm dimensions of the NPs compressed at a \u003cem\u003eT\u003c/em\u003e\u003csub\u003ecm\u003c/sub\u003e of 128\u0026deg;C (Extended Data Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ea).\u003c/p\u003e \u003cp\u003eAs shown in Extended Data Fig.\u0026nbsp;7, the isothermally crystallized G1-3 powders pressed at 140\u0026deg;C also show highly ordered anisotropic meso-morphology both in the in-plane and through-plane directions. Long-period structures with a wide range of domain dimensions of 71.7\u0026ndash;88.8 nm, 62.8\u0026ndash;80.8 nm and 53.7\u0026ndash;64.3 nm are observed for the G1-3 disks, which is in accordance with their high \u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e values (Extended Data Fig.\u0026nbsp;8a).\u003c/p\u003e \u003cp\u003eIn comparison, the microscale 2D-ordered nanofibrils of 22.9\u0026thinsp;\u0026plusmn;\u0026thinsp;4.1 nm and macroscale randomly 3D-organized nanofibrils self-assembled by pressing at a \u003cem\u003eT\u003c/em\u003e\u003csub\u003ecm\u003c/sub\u003e of 145\u0026ndash;220\u0026deg;C followed by isothermal crystallization at a \u003cem\u003eT\u003c/em\u003e\u003csub\u003ec\u003c/sub\u003e of 128\u0026deg;C for 3\u0026ndash;10 h. The macroscale randomly 3D-distributed nanofibrils of 23.9\u0026thinsp;\u0026plusmn;\u0026thinsp;4.7 nm self-organized by quenching from the \u003cem\u003eT\u003c/em\u003e\u003csub\u003ecm\u003c/sub\u003e of 190\u0026deg;C. The long nanofibrils and their aggregates are 3D inter-crosslinked by short nanofibers of 71\u0026ndash;101 nm and knots. The tensile properties of unoriented G4 sheets strongly depend on the entangled networks at elevated \u003cem\u003eT\u003c/em\u003e\u003csub\u003ecm\u003c/sub\u003e and microscale 2D-ordered and randomly 3D-organized nanofibrillar superstructures (Extended Data Fig.\u0026nbsp;6c and Fig. \u003cspan refid=\"MOESM1\" class=\"InternalRef\"\u003eS1\u003c/span\u003e). The disentangled nanofibrillar superstructure has the lowest ductility (Fig. \u003cspan refid=\"MOESM1\" class=\"InternalRef\"\u003eS1\u003c/span\u003e-2).\u003c/p\u003e \u003cp\u003eAs shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ec, the thick G4 bars have the highest impact strength of 202\u0026thinsp;\u0026plusmn;\u0026thinsp;16.6 KJ/m\u003csup\u003e2\u003c/sup\u003e among all the synthetic polymers. The G5 bars with a high content of nanofibrils exhibit a moderate impact resistance of 101\u0026thinsp;\u0026plusmn;\u0026thinsp;12 KJ/m\u003csup\u003e2\u003c/sup\u003e, which is higher than the 83.3\u0026thinsp;\u0026plusmn;\u0026thinsp;1.2 KJ/m\u003csup\u003e2\u003c/sup\u003e of the G3 sheet and the extremely low impact strength of iPP. This result indicates that there is a threshold concentration of hierarchical meso-particles or meso-/nanofibrillar networks that strongly influence the orientation of nanofibrillar aggregates and absorb energy during impact deformation (Extended Data Fig.\u0026nbsp;6a/8b). For the sheets from the G5 powders with a relatively high concentration of intrinsic nanofibrils, the nanofibrillar networks cannot be easily deformed, the networks may break, and only discontinuous bands of deformation are generated. However, the impact resistance of the G5 sheets is still high enough because of their homogenized nanofibrillar structures.\u003c/p\u003e \u003c/div\u003e\n\u003ch3\u003eTransition from 2D-Order Meso-fibrils to Disorder 3D Granules\u003c/h3\u003e\n\u003cp\u003eHierarchical order-to-disorder melt transformations during vacuum annealing are shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ea. The transformation from intra-crosslinked 2D meso-/nanofibrils to disordered 3D granules was observed after the annealing of the NPs at 250\u0026deg;C for 6\u0026ndash;20 h under vacuum (Extended Data Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ea). The reprocessed and isothermal crystallized G2 sheets have \u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e values of 135.4\u0026deg;C and 130.8\u0026deg;C, respectively, according to the first and second heating curves obtained via DSC. The quenched G2 sheet exhibited interesting ultrahigh tensile ductility, with an average tensile strain at break of 1532\u0026thinsp;\u0026plusmn;\u0026thinsp;95%, a yield strength of 29.38\u0026thinsp;\u0026plusmn;\u0026thinsp;0.72 MPa and a low tensile strength at break of 31.73\u0026thinsp;\u0026plusmn;\u0026thinsp;1.85 MPa. Isothermal crystallization does not significantly influence the tensile properties of G2 (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003eb). Unexpectedly, the crystallinity of these sheets is as high as 70%. The G2 sheet shows completely distinct tensile deformation behaviors, with strong tensile yielding, long neck development and much lower tensile strain hardening than its inter-crosslinked counterparts. In contrast, the disentangled disordered G2 sheet shows tensile brittleness at a deformation temperature of 120\u0026deg;C (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ec).\u003c/p\u003e \u003cp\u003eIn comparison, the vacuum-annealed low-molecular-weight G3 powders tend to coalesce quickly and undergo chain entanglement, maintaining their order melt (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ea). The vacuum-annealed high-molecular-weight G6 powders containing high contents of intrinsic nascent nanofibrils do not compact each other (Extended Data Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003eb). After vacuum annealing, the compressed G3 sheets exhibited high tensile ductility after room-temperature and high-temperature tensile deformation (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003eb-c and Extended Data Fig.\u0026nbsp;5f). The G3 sheets show much lower melt drawability, a low melt strength of 0.014 MPa, and a dual-shaped recovery of 48%. After static re-entangling at 220\u0026deg;C for 5 min, the G3 sample has a higher melt strength of 0.055 MPa and a dual-shaped recovery of 72.7% (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ed and Extended Data Fig.\u0026nbsp;5g).\u003c/p\u003e \u003cp\u003eAs shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ea, isothermal crystallization at 128\u0026deg;C induced a transition from an ordered melt to a higher-order condensed solid-state in the G3 sheet. In contrast, isothermal crystallization at 116\u0026ndash;124\u0026deg;C does not result in a transition from a disordered melt to a condensed ordered solid-state.\u003c/p\u003e\n\u003ch3\u003eMeso-Superstructures from Inter-Crosslinked Order Melts\u003c/h3\u003e\n\u003cp\u003eAnnealing at 142\u0026ndash;300\u0026deg;C under pressure suppresses the order-to-disorder melt transition and chain explosion. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ea and Extended Data Fig.\u0026nbsp;7/8c, the isothermal crystallized and quenched G1-2 sheets, even when pressed at a \u003cem\u003eT\u003c/em\u003e\u003csub\u003ecm\u003c/sub\u003e of 300\u0026deg;C, maintain the ordered melt state. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003eb, the isothermal crystallized G2 sheet from 300\u0026deg;C has an average high tensile strength of 75.6\u0026thinsp;\u0026plusmn;\u0026thinsp;2.4 MPa, similar to the 76.6\u0026thinsp;\u0026plusmn;\u0026thinsp;3.1 MPa of the G1 sheet. The isothermal crystallized and quenched G2 sheets from 142\u0026deg;C show higher tensile strengths of 63.1\u0026thinsp;\u0026plusmn;\u0026thinsp;3.2 and 49.5\u0026thinsp;\u0026plusmn;\u0026thinsp;25 MPa than the 48.8\u0026thinsp;\u0026plusmn;\u0026thinsp;2.1 and 42.4\u0026thinsp;\u0026plusmn;\u0026thinsp;1.8 MPa of the G1 sheets (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003eb, Extended Data Fig.\u0026nbsp;8d and Table \u003cspan refid=\"MOESM1\" class=\"InternalRef\"\u003eS1\u003c/span\u003e). The G2 sheets compressed at 142\u0026deg;C also exhibited much greater tensile drawability at the testing temperature of 120\u0026deg;C, with strong strain-hardening properties (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ec). The inter-crosslinked G2 sheet has a high melt strength of 0.367 MPa and melt\u0026ndash;tensile ductility (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ed). These results show that the tensile properties of G2 sheets from 2D meso-fibrils do not depend on the processing temperature or crystallinity due to hierarchical meso-fibrillar and nanofibrillar networks. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ed, the inter-crosslinked G2 sheets show excellent shape recovery of 116% (Extended Data Fig.\u0026nbsp;5g). Moreover, the isothermal crystallized G2 sheet from 300\u0026deg;C possesses unusual optical transparency and unique selective transmission of visible light. The isothermally crystallized and quenched G2 sheets with thicknesses of 1.1/1.53 \u0026micro;m show high transmittances of 79.8/84.2% and 31/34% for near-infrared and ultraviolet light, respectively, at wavelengths of 1600 and 350 nm (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ee). This unique optical transparency was observed only in OBC with a layered mesophase morphology and rod-like lamellar crystals\u003csup\u003e31\u003c/sup\u003e.\u003c/p\u003e \u003cp\u003eThe qualitative 2D-USAXS patterns in Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ee-f present hierarchical ordered anisotropic superstructures in both the in-plane and through-plane directions. The quantitative 1D-USAXS curves show that the isothermally crystallized G2 sheet from 300\u0026deg;C has meso-scale superstructures with long periodic dimensions of 49.6\u0026ndash;235.6 nm. The quenched G2 sheet from 300\u0026deg;C has a nanoscale superstructure with long periodic dimensions of 36.6\u0026ndash;49.5 nm, which is similar to that of the NPs. The isothermally crystallized G1-2 sheets from 142\u0026deg;C have meso-scale superstructures with long periodic dimensions of 88.8\u0026ndash;137.7 nm and 44.6\u0026ndash;182.6 nm (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ef and Extended Data Fig.\u0026nbsp;7), respectively. The G2 sheets with 2D-oriented nanofibrils and the G4 sheets with 2D-ordered nanofibrils have no strongly oriented orthogonal crystal-stems (Extended Data Fig.\u0026nbsp;8e and Fig. S3). The orientation parameter of the crystal planes (A\u003csub\u003e[200]\u003c/sub\u003e/A\u003csub\u003e[110]\u003c/sub\u003e) is defined as the peak area ratio of the [200] and [110] reflections of the orthogonal crystals in E-PE. The A\u003csub\u003e[200]\u003c/sub\u003e/A\u003csub\u003e[110]\u003c/sub\u003e values of the 2D-oriented G2, 2D-ordered G4 and 3D-random G4 nanofibrils are 0.5, 0.29 and 0.09, respectively, which are much lower than those of high-strength tapes\u003csup\u003e7\u003c/sup\u003e.\u003c/p\u003e\n\u003ch3\u003eEplastomers via Dynamic Nanofibrillar Networks\u003c/h3\u003e\n\u003cp\u003eAs shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ea and Extended Data Fig.\u0026nbsp;9, the eplastomers, composed of hierarchical entangled 5 wt% G1 nanofibrillar networks in the polyolefin elastomer (POE) matrix, possess self-reinforced soft\u0026ndash;plastic properties after extrusion and compression molding below \u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e. The nanofibrillar composite exhibited a high yield strength of 22.9\u0026thinsp;\u0026plusmn;\u0026thinsp;4.35 MPa, which was much greater than the value of 3.61\u0026thinsp;\u0026plusmn;\u0026thinsp;0.1 MPa for POE\u003csub\u003eCM\u003c/sub\u003e at a fixed strain of 50%. The extruded and injection-molded eplastomer (E2\u0026thinsp;+\u0026thinsp;IM) at a \u003cem\u003eT\u003c/em\u003e\u003csub\u003ep\u003c/sub\u003e of 160\u0026deg;C also behaves as a soft plastic with a high tensile yield strength of 13.93\u0026thinsp;\u0026plusmn;\u0026thinsp;0.34 MPa due to the shear-induced orientation of the nanofibrils, which is much greater than the 5.35\u0026thinsp;\u0026plusmn;\u0026thinsp;0.09 MPa of POE\u003csub\u003eIM\u003c/sub\u003e at a fixed strain of 150%. After compression mold at a \u003cem\u003eT\u003c/em\u003e\u003csub\u003ep\u003c/sub\u003e of 160\u0026deg;C, the epalstomer exhibited elastomeric properties due to the transformation of 2D-ordered nanofibrillar networks to disentangled disordered 3D meso-/nanoparticles. The meso-particle composites have a low tensile strength of 4.23\u0026thinsp;\u0026plusmn;\u0026thinsp;0.08 MPa at a fixed strain of 50%, a high tensile strength at break of 22.24\u0026thinsp;\u0026plusmn;\u0026thinsp;0.81 MPa and a tensile strain of 1398\u0026thinsp;\u0026plusmn;\u0026thinsp;40%. The re-extruded and re-compressed nanofibrillar composites (E3\u0026thinsp;+\u0026thinsp;CM3) also exhibited soft\u0026ndash;plastic properties (Fig. S4).\u003c/p\u003e \u003cp\u003eAs shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003eb-e and Extended Data Fig.\u0026nbsp;9, the nanofibrillar crystals have high \u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e values ranging from 132.3\u0026deg;C to 158\u0026deg;C (Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003eb). The DSC curves show that the non-crosslinked high-order G1 meso-particles can be completely dispersed in the POE matrix (Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003eb). At temperatures of 115\u0026deg;C and 125\u0026deg;C below the \u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e of the nanofibrils, the compressed nanofibrillar composites have a much higher complex viscosity of 3.22\u0026thinsp;~\u0026thinsp;4.25 \u0026times; 10\u003csup\u003e4\u003c/sup\u003e Pa\u0026middot;s than does the 7.65 \u0026times; 10\u003csup\u003e3\u003c/sup\u003e pure POE at a fixed angular frequency of 0.1 rad/s (Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ec). The nanofibrillar networks have a 1.97/2.89 times greater diffraction peak intensity for the [110] phase than for the disordered meso-particles, regardless of the compression or injection mold processes (Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ed). The qualitative 2D-USAXS patterns show that the eplastomers are different from the SMCs of the HDPE/E-PE and OBC/E-PE composites reinforced with shish-kebab crystals\u003csup\u003e33\u003c/sup\u003e. Figure\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ee shows the toluene-etched morphologies of the compressed and injection-molded eplastomers. Homogeneously dispersed nanofibrillar networks are observed. The nanofibrils had an average diameter of 51.1\u0026thinsp;\u0026plusmn;\u0026thinsp;23.8 nm. Interestingly, the width of these nanofibrils is equal to that of traditional lamellar crystals and the radius of gyration\u003csup\u003e40\u003c/sup\u003e, but the length of the nanofibrils is unexpectedly greater than 10\u0026ndash;50 \u0026micro;m.\u003c/p\u003e \u003cp\u003eThe disordered meso-/microparticles in the POE1-2 matrix have a low melting peak at 127\u0026deg;C (Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003eb), which is in line with the low \u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e of 126.4\u0026deg;C for high-temperature annealed disordered G1-NPs. The meso-particle composites had slightly greater complex viscosities than did pure POE at low angular frequencies (Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ec). This result shows that the eplastomers are re-processible and can be melt-mechanically recycled into high-strength soft plastics. The compressed meso-particle composites contain separated and less aggregated meso-granules with average diameters of 640\u0026ndash;881 nm and microgranules with average diameters of 1.24\u0026ndash;1.95 \u0026micro;m (Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ee). The abovementioned results prove that the non-crosslinked meso-particles of G1 have potential applications in reinforcing elastomers, rubber and soft plastics and replacing inorganic fillers.\u003c/p\u003e \u003cp\u003eIn summary, polymerization-induced nano/meso-ordering determines the order of melt below \u003cem\u003eT\u003c/em\u003e\u003csub\u003eO\u0026minus;D\u003c/sub\u003e, order-to-disorder melt transformation at \u003cem\u003eT\u003c/em\u003e\u003csub\u003eO\u0026minus;D\u003c/sub\u003e, chain explosion and re-entanglements at \u003cem\u003eT\u003c/em\u003e\u003csub\u003eO\u0026minus;D\u003c/sub\u003e and meso-order condensed crystalline superstructures from the ordered melt. The order-to-disorder melt transition also significantly depends on the morphology, molecular weight and processing pressure of the NPs. Meso-order condensed superstructures result in unique optical properties. The use of heterogeneous entangled and inter-crosslinked networks results in excellent dual-shaped memory properties. However, homogeneous re-entangled networks lead to low melt ductility.\u003c/p\u003e"},{"header":"Declarations","content":"\u003cp\u003e\u003cstrong\u003eData availability\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eThe data that support the findings of this study are presented in the paper and/or the Supplementary Information and are available from the corresponding authors on request.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eAcknowledgments\u003c/strong\u003e\u0026nbsp;\u003c/p\u003e\n\u003cp\u003eThe founding of this research is supported by the National Natural Science Foundation of China (Grant No. U19B6001) and the National Key R\u0026amp;D Plan of China (Grant No. 2020YFC1107003).\u0026nbsp;We thank the beamline BL10U1 station of the \u003cem\u003eShanghai Synchrotron Radiation Facility\u0026nbsp;\u003c/em\u003efor 2D-USAXS measurements. We also thank Yonglong Wu, Xiaoliang Wang and Shunbing Wang for high-temperature solid-state tensile stress–strain testing.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eAuthor contributions\u003c/strong\u003e\u0026nbsp;\u003c/p\u003e\n\u003cp\u003eA.M. developed the original idea, designed and conducted all the experiments, developed the methods, and calculated and analyzed all the results; K.C. provided processing equipment and performed the rheological testing; F.T. performed the 2D-USAXS experiments. A.M. wrote the initial manuscript, revised it, and edited various subsequent versions.\u0026nbsp;\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eCompeting interests\u003c/strong\u003e\u0026nbsp;\u003c/p\u003e\n\u003cp\u003eThe authors declare that they have no competing interests.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eAdditional information\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eSupplementary information\u003c/strong\u003e\u0026nbsp;\u003c/p\u003e\n\u003cp\u003eThe online version contains supplementary material available at https://doi.org/X/X.\u0026nbsp;\u003c/p\u003e\n\u003cp\u003eCorrespondence and requests for materials should be addressed to Aizezi Maimaitiming.\u0026nbsp;\u003c/p\u003e\n\u003cp\u003e\u003cb\u003eOnline Content\u003c/b\u003e\u003c/p\u003e\n\u003cp\u003eAny methods, additional references, Nature Portfolio reporting summaries, source data, extended data, supplementary information, acknowledgments, peer review information; details of author contributions and competing interests; and statements of data and code availability are available at https://doi.org/XXX.\u003c/p\u003e"},{"header":"References","content":"\u003col\u003e\u003cli\u003e\u003cspan\u003eMarie, F. 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Nanotechnol. 5, 251\u0026ndash;255 (2010).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003ePan, X., Shen, L., Albertus, P. H. J. S. \u0026amp; Cees, W. M. B. Transparent, high-thermal-conductivity ultradraw polyethylene/graphene nanocomposite films. Adv. Mater. 31, 1904348 (2019).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eHiroki, U. et al. Continuous film processing from ultrahigh-molecular-weight polyethylene reactor powder and mechanical property development by melt drawing. Ind. Eng. Chem. Res. 45, 7801\u0026ndash;7806 (2006).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eLi, R. et al. Scalable production of critically thin polyethylene films via multistep stretching. Nat. Chem. Eng. 1, 702\u0026ndash;709 (2024).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eRob, C. P. V. et al. Fast, light-responsive, metal-like polymer actuators generating high stresses at low strain. \u003cem\u003eMatter\u003c/em\u003e 2, 1522\u0026thinsp;\u0026ndash;\u0026thinsp;1534 (2020).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eLiu, X. et al. Transparent ultrahigh-molecular-weight polyethylene/MXene films with efficient UV-absorption for thermal management. Nat. Commun. 15, 3076 (2024).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003e\u0026Aacute;kos, K., Tam\u0026aacute;s, B. \u0026amp; J\u0026oacute;zsef, K. K. Self-reinforced polymeric materials: A review. Prog. Polym. Sci. 35, 1288\u0026ndash;1310 (2010).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eYao, Q., Lisa, D. F., Madhusudhan, R. P. \u0026amp; Kevin, L. S. A review of the fabrication methods and mechanical behavior of continuous thermoplastic polymer fiber-thermoplastic polymer matrix composites. Polym. Compos. 44, 694\u0026ndash;733 (2023).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eWinfried, B., Bob, F. \u0026amp; Pieter, J. L. Chapter 2: High-performance fibers and tapes based on PP and PE. Adv. Ind. Eng. Polym. Res. 5, 60\u0026ndash;69 (2022).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eMasaki, K. \u0026amp; Dai, F. Preparation of ultrahigh-molecular-weight polyethylene fibers by combination of melt-spinning and melt-drawing. Mater. Today Commun. 23, 100864 (2020).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eStefano, O., Elena, F., Adele, F., Viscardo, M. \u0026amp; Roger, S. P. SAXS investigations on uniaxially drawn fibers obtained from polyethylene reactor powder. \u003cem\u003eMacromolecules\u003c/em\u003e 29, 3292\u0026thinsp;\u0026ndash;\u0026thinsp;3299 (1996).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eThein, K., Kenichi, F., Myung, H. C., Takeshi, K. \u0026amp; Jar, S. L. Melting and crystallization of gelation-crystallized ultrahigh-molecular-weight polyethylene. Macromolecules 22, 2238\u0026ndash;2244 (1989).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eHiroki, U., Mitsuhiro, N., Tetsuo, K., Anagnostis, E. Z. \u0026amp; Roger, S. P. Melt drawability of ultrahigh molecular weight polyethylene. Macromolecules 32, 2761\u0026ndash;2769 (1999).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eDeplancke, T. et al. Diffusion versus cocrystallization of very long polymer chains at interfaces: experimental study of sintering of UHMWPE nascent powder. \u003cem\u003eMacromolecules\u003c/em\u003e 47, 197\u0026thinsp;\u0026ndash;\u0026thinsp;207 (2014).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eHiroki, U., Tetsuo, K., Akiyoshi, K. \u0026amp; Syozo, M. Real-time X-ray diffraction study on two-stage drawing of ultra-high molecular weight polyethylene reactor powder above the static melting temperature. \u003cem\u003eMacromolecules\u003c/em\u003e 29, 1540\u0026thinsp;\u0026ndash;\u0026thinsp;1547 (1996).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eMasaki, K. et al. Phase Transitions during heating of melt-drawn ultrahigh molecular weight polyethylenes having different molecular characteristics. J. Phys. Chem. B 112, 5311\u0026ndash;5316 (2008).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eFeng, Y., Tong, W., Ming, X. \u0026amp; Ya, C. Deformation and pore formation mechanism of β nucleated polypropylene with different supermolecular structures. Eur. Polym. J. 91, 134\u0026ndash;148 (2017).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eEdward, B. T. et al. Self-assembled highly ordered acid layers in precisely sulfonated polyethylene produce efficient proton transport. Nat. Mater. 17, 725\u0026ndash;731 (2018).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eLu, Y., Christina, R., Stefan, M. \u0026amp; Karen, I. W. Gyroid and other ordered morphologies in single-ion conducting polymers and their impact on ion conductivity. J. Am. Chem. Soc. 142, 857\u0026ndash;866 (2020).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eAizezi, M. et al. Thermal insulative 2D-oriented bimodal micro-foams and high strength photonic thermoplastic elastomers by assembling olefin multiblock copolymers. Compos. B. Eng. 246, 110283 (2022).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eJinwoo, O. et al. Shear-Rolling process for unidirectionally and perpendicularly oriented sub-10-nm block copolymer patterns on the 4 in scale. ACS Nano 15, 8549\u0026ndash;8558 (2021).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eTimo, H., Carl, G. S., Patrizia, P., Daniel, H. \u0026amp; Rolf, M. Self-Reinforcement via 1D nanostructure formation during melt blending of thermoplastics and thermoplastic elastomers with nanophase-separated UHMWPE/HDPE wax reactor blends. ACS Appl. Polym. Mater. 3, 3455\u0026ndash;3464 (2021).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eWang, X., Li, H. \u0026amp; Jin. R. Synchronous toughening and reinforcing of polypropylene with ultrahigh-molecular-weight polyethylene via melt blending: mechanical properties, morphology, and rheology. J. Appl. Polym. Sci. 100, 3498\u0026ndash;3509 (2006).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eJauffr\u0026egrave;s, D., Lame, O., Vigier, G. \u0026amp; Dor\u0026eacute;, F. How nascent structure of semicrystalline polymer powders enhances bulk mechanical properties. \u003cem\u003eMacromolecules\u003c/em\u003e 41, 9793\u0026thinsp;\u0026ndash;\u0026thinsp;9801 (2008).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eRamin, H. et al. Formation of UHMWPE nanofibers during solid-state deformation. Nanomaterials 12, 3825 (2022).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eDe Gennes, P.-G. Explosion \u0026agrave; la fusion (explosion upon melting). C. R. Acad. Sci., Ser. IIb: Mec., Phys., Astron. 321, 363\u0026ndash;365 (1995).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eTiana, D., Olivier, L., Fran\u0026ccedil;ois, R., Roland, S. \u0026amp; G\u0026eacute;rard, V. Mechanisms of chain reentanglement during the sintering of UHMWPE nascent powder: effect of molecular weight. \u003cem\u003eMacromolecules\u003c/em\u003e 48, 5328\u0026thinsp;\u0026ndash;\u0026thinsp;5338 (2015).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eAmanda, G. M., Paul, J. D., Jeff, S. F., Ronald, L. J. \u0026amp; Chad, R. S. Measuring tie chains and trapped entanglements in semicrystalline polymers. \u003cem\u003eMacromolecules\u003c/em\u003e 53, 5614\u0026thinsp;\u0026ndash;\u0026thinsp;5626 (2020).\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eJohn, D. H. \u0026amp; Robert, L. M. Kinetic of crystallization from the melt and chain folding in polyethylene fractions revisited: theory and experiment. \u003cem\u003ePolymer\u003c/em\u003e 38, 3151\u0026thinsp;\u0026ndash;\u0026thinsp;3212 (1997).\u003c/span\u003e\u003c/li\u003e\u003c/ol\u003e"},{"header":"Methods","content":"\u003cp\u003e\u003cstrong\u003eMaterials\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eCommercial nascent G1\u0026ndash;6 powders were purchased from Ticona Co., with the following grade names: GUR\u003csup\u003e\u0026reg;\u003c/sup\u003e2024 (G1), GUR\u003csup\u003e\u0026reg;\u003c/sup\u003e4120 (G2), GUR\u003csup\u003e\u0026reg;\u003c/sup\u003e2122 (G4) and GUR\u003csup\u003e\u0026reg;\u003c/sup\u003e4170 (G5). The viscosity-average molecular weights (\u003cem\u003eM\u003c/em\u003e\u003csub\u003e\u0026eta;\u003c/sub\u003e) of G1\u0026minus;5 are 5.4E6, 4.2E6, 6.9E5, 4.7E6 and 1.02E7\u0026nbsp;g/mol (Margolies\u0026rsquo; equation). The polyolefin elastomers (POE1\u0026minus;2) were purchased from Dow Chemical Co. The chemical structures of POE1 and iPP are shown in ref. 41, and the grade names of POE2 and iPP are Engage\u003csup\u003eTM\u003c/sup\u003e 8440 and PPT30S, respectively. The melt flow indices of POE1\u0026minus;2 are 5 and 1 g/10 min (190 \u0026deg;[email protected] kg), respectively. The antioxidant 1010 was purchased from BASF Co.\u003c/p\u003e\n\u003cp\u003e\u003cb\u003eCharacterization and Testing\u003c/b\u003e\u003c/p\u003e\u003cp\u003e\u003cb\u003eHierarchical Superstructures\u003c/b\u003e:\u003c/p\u003e\u003cp\u003eThe morphologies of the nascent powders, nanofibrils, meso-particles and fractured surfaces of the sheets were characterized by scanning electron microscopy (SEM, 2\u0026thinsp;\u0026minus;\u0026thinsp;10 keV, 5\u0026thinsp;\u0026minus;\u0026thinsp;10 mA, HITACHI Regulus 8100). Before characterization, the dumbbell-shaped samples and sheets were cryo-fractured in liquid nitrogen. The surfaces of nascent powders and cryo-fractured sheets were doped with platinum via a HITACHI MCI000. Before the characterization of morphologies, the cryo-fractured surfaces of eplastomers (95POE1\u0026thinsp;\u0026minus;\u0026thinsp;2/5G1) emerged in xylene at 80\u0026deg;C for 2\u0026ndash;6 h to etch the POE matrix phase. The samples were subsequently dried at 60\u0026deg;C for 6\u0026ndash;24 h. Quantitative analysis was performed via Image-Pro Plus 6.0 software.\u003c/p\u003e\u003cp\u003eStatic transformations from condensed ordering states to order-to-disorder melts of nascent G1\u0026ndash;3 powders are tested via differential scanning calorimetry (DSC2500, TA Instrument). Pieces of 4\u0026ndash;8 mg of powder were heated from 25\u0026deg;C to an annealing temperature (\u003cem\u003eT\u003c/em\u003e\u003csub\u003ea\u003c/sub\u003e) of 140\u0026ndash;400\u0026deg;C at a heating rate of 10\u0026deg;C/min, isothermal for 2\u0026ndash;5 h, subsequently cooled to 40\u0026deg;C at a cooling rate of 10\u0026deg;C/min, isothermal for 3 min, again heated to an annealing temperature of 160\u0026deg;C, isothermal for 5 min, and cooled to 25\u0026ndash;40\u0026deg;C at a cooling rate of 10\u0026deg;C/min. The melting temperature (\u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e) and crystallization temperature (\u003cem\u003eT\u003c/em\u003e\u003csub\u003ec\u003c/sub\u003e) are obtained from the heating and cooling curves, respectively.\u003c/p\u003e\u003cp\u003eThe order-to-disorder melt transformation temperature (T\u003csub\u003eO\u0026minus;D\u003c/sub\u003e) is characterized via a sequential annealing and a subsequent non-isothermal crystallization and melting (SA-SC-M) procedure via differential scanning calorimetry (DSC2500, TA Instrument). A 4\u0026ndash;8 mg sample of powder or sheet was heated from 25\u0026ndash;40\u0026deg;C to an annealing temperature (\u003cem\u003eT\u003c/em\u003e\u003csub\u003ea\u003c/sub\u003e) at a heating rate of 10\u0026deg;C/min, isothermal for 1 h, subsequently cooled to 25/40\u0026deg;C at a cooling rate of 10\u0026deg;C/min, isothermal for 3 min, again heated to the annealing temperature of \u003cem\u003eT\u003c/em\u003e\u003csub\u003ea\u003c/sub\u003e + 50\u0026deg;C, and the subsequent crystallization procedure and sequential annealing procedure were repeated under the same conditions. \u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e and \u003cem\u003eT\u003c/em\u003e\u003csub\u003ec\u003c/sub\u003e are obtained from the heating and cooling curves, respectively.\u003c/p\u003e\u003cp\u003eCrystal melting, chain entanglements above \u003cem\u003eT\u003c/em\u003e\u003csub\u003em\u003c/sub\u003e, chain explosion, re-entanglements and relaxation above \u003cem\u003eT\u003c/em\u003e\u003csub\u003eO\u0026minus;D\u003c/sub\u003e are measured in situ via a torque rheometer (HAAKE PolyLab OS Rheo-Drive 4, Thermo-Scientific Company). The rotation speed is 30\u0026ndash;60 rpm. The shearing temperature increases continuously from 130\u0026deg;C to 250\u0026deg;C at intervals of 5\u0026deg;C for 2\u0026ndash;4 min. The data acquisition of the torque values is manual from the rheometer software (Extended Data Fig.\u0026nbsp;10a).\u003c/p\u003e\u003cp\u003eThe condensed ordering phases of the pressed sheets are analyzed via two-dimensional ultrasmall-angle X-ray scattering experiments (2D-USAXS, \u003cem\u003eλ\u003c/em\u003e\u0026thinsp;=\u0026thinsp;0.124 nm, 10 keV), which are performed at the beamline BL10U1 station of the \u003cem\u003eShanghai Synchrotron Radiation Facility\u003c/em\u003e. The exposure time was 2\u0026ndash;20 s. A Dectris Eiger 4 M with a pixel size of 75 \u0026times; 75 \u0026micro;m was used to collect the signals. The distance between the sample and the detector was 27600 mm, which was calibrated via a bovine tendon standard sample. The scattering data are background corrected according to a general procedure.\u003c/p\u003e\u003cp\u003eThe crystal stems and lamellar crystals were analyzed via two-dimensional small-/wide-angle X-ray scattering experiments (2D-SAXS/WAXS, \u003cem\u003eλ\u003c/em\u003e\u0026thinsp;=\u0026thinsp;0.124 nm, 10 keV), which were performed at the \u003cem\u003eShanghai Synchrotron Radiation Facility\u003c/em\u003e. The exposure time was 10 s. A Pilatus 2 M detector (Dectris Swiss) with a pixel size of 172 \u0026times; 172 \u0026micro;m was used to collect the signals.\u003c/p\u003e\u003cp\u003eCrystal stems and orientation were also characterized via X-ray powder diffraction (XRD). XRD patterns were obtained with CuKa (\u003cem\u003eλ\u003c/em\u003e\u0026thinsp;=\u0026thinsp;0.15418 nm) radiation via an X-ray diffractometer (X\u0026rsquo;Pert Powder, PAN analytical), performing a continuous scan at diffraction angles of 2\u003cem\u003eθ\u003c/em\u003e\u0026thinsp;=\u0026thinsp;3\u0026ndash;50\u0026deg; at 40 mV and 40 mA.\u003c/p\u003e\u003cp\u003eThe quantitative orientation ratio of the [200] and [110] crystal phases of the orthogonal crystal stems of E-PE is calculated as follows\u003csup\u003e7\u003c/sup\u003e:\u003cdiv id=\"Equa\" class=\"Equation\"\u003e\u003cdiv format=\"TEX\" class=\"mathdisplay\" id=\"FileID_Equa\" name=\"EquationSource\"\u003e\n$$\\:\\frac{{A}_{\\left[200\\right]}}{{A}_{\\left[110\\right]}}=100\\:\\times\\:\\:\\frac{{A}_{\\left[200\\right]}}{{A}_{\\left[110\\right]}}$$\u003c/div\u003e\u003c/div\u003e\u003c/p\u003e\u003cp\u003ewhere \u003cem\u003eA\u003c/em\u003e\u003csub\u003e[200]\u003c/sub\u003e and \u003cem\u003eA\u003c/em\u003e\u003csub\u003e[110]\u003c/sub\u003e are the intensities of the [200] and [110] peaks of the orthogonal crystal stems at 2\u003cem\u003eθ\u003c/em\u003e\u0026thinsp;=\u0026thinsp;21.5\u0026deg;/21.7\u0026deg; and 23.9\u0026deg;/24.1\u0026deg;, respectively, in the XRD profiles.\u003c/p\u003e\u003cp\u003eThe crystallinity is calculated as follows:\u003cdiv id=\"Equb\" class=\"Equation\"\u003e\u003cdiv format=\"TEX\" class=\"mathdisplay\" id=\"FileID_Equb\" name=\"EquationSource\"\u003e\n$$\\:{\\chi\\:}_{c}=\\frac{{\\varDelta\\:H\\times\\:100}_{\\:}}{{\\varDelta\\:H}_{PE}^{*}}$$\u003c/div\u003e\u003c/div\u003e\u003c/p\u003e\u003cp\u003ewhere ∆\u003cem\u003eH\u003c/em\u003e is the melting enthalpy in the heating DSC curves of the samples. ∆\u003cem\u003eH\u003c/em\u003e\u003csub\u003ePE\u003c/sub\u003e* is the melting enthalpy of totally crystallized polyethylene with a value of 290 J/g\u003csup\u003e41\u003c/sup\u003e.\u003c/p\u003e\u003cp\u003eSolid-state tensile stress\u0026ndash;strain tests of dumbbell- and bar-shaped samples cut from the sheets were performed according to the standard methods of ISO527.2-5B/50, ISO20573-TypeCP/50 or ISO527.2-5A/50 (Instron 5966 and Instron 34TM-10). The average values of at least 3\u0026ndash;6 tests with high reproducibility were recorded in the tensile strain\u0026ndash;stress curves.\u003c/p\u003e\u003cp\u003eHigh-temperature solid\u0026ndash;state tensile stress\u0026ndash;strain tests of dumbbell-shaped samples cut from the sheets were performed at a testing temperature of 120\u0026deg;C according to the standard methods of ISO527.2-5B/5 (CMT4000, Shenzhen SAS Test Equipment Co. Ltd.). The average values of at least 2\u0026ndash;3 tests with high reproducibility were recorded in the tensile stress\u0026ndash;strain curves.\u003c/p\u003e\u003cp\u003eThe melt-state tensile stress\u0026ndash;strain properties of the UHMWPE sheets were measured via an ARES-G2 rheometer (TA Instruments). A sheet sample with a size of 35 \u0026times; 2 \u0026times; 2 mm was cut from the sheets and compressed and molded at distinct temperatures. The testing temperature was 150\u0026deg;C. The soaking time for every annealing temperature prior to testing was 120 s. The constant linear rate was 2.5\u0026ndash;6 mm/min.\u003c/p\u003e\u003cp\u003eThe dual shape memory effects of the UHMWPE sheets were measured via an ARES-G2 rheometer (TA Instruments). A disentangled G3 sheet and an inter-crosslinked G2 sheet with a size of 35 \u0026times; 2 \u0026times; 2 mm were cut from the compression-molded sheets. The samples are melt-drawn at a rate of 2.5 mm/min to a strain of 160% at 150\u0026deg;C, and the temporary shape is fixed under stress after cooling to 40\u0026deg;C. The samples were subsequently recovered into their initial shapes by heating to 150/220\u0026deg;C. The G4 samples were re-entangled at 220\u0026deg;C for 5 min before testing. The shape recovery was calculated according to the standard method\u003csup\u003e41\u003c/sup\u003e.\u003c/p\u003e\u003cp\u003eBi-notched Charpy impact testing for UHMWPE sheets was performed using standard bar samples with a size of 100 \u0026times; 10 \u0026times; 4 mm, and two notches with 45\u0026deg; \u0026plusmn; 1\u0026deg; and R\u0026thinsp;=\u0026thinsp;0.25\u0026thinsp;\u0026plusmn;\u0026thinsp;0.05 mm were machined. At least 6 notched bar samples were tested at an impact energy of 50 J (KD-XJJD-50).\u003c/p\u003e\u003cp\u003eNotched Izod impact testing for entangled G3 and iPP bars was performed using standard bar samples with a size of 100 \u0026times; 10 \u0026times; 4 mm, and one notch with 45\u0026deg; \u0026plusmn; 1\u0026deg; and R\u0026thinsp;=\u0026thinsp;0.25\u0026thinsp;\u0026plusmn;\u0026thinsp;0.05 mm was machined. At least 4\u0026ndash;5 bar samples were tested at an impact energy of 5.5 J, CESAT9050-5.5.\u003c/p\u003e\u003cp\u003eThe rheological properties of the UHMWPE and 95POE1/5G1 sheets were measured via an Anton-Paar MCR 302 rheometer. A disk sample with a diameter of 25 mm and a thickness of 1\u0026ndash;2 mm is used. The strain is 5%, and the gap is 1\u0026ndash;2 mm. The angular frequency is 0.1\u0026ndash;100 rad/s, and the testing temperatures are 115\u0026ndash;220\u0026deg;C. The soaking time for every annealing temperature prior to testing was 2\u0026ndash;10 min.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003ePreparation and Processing of UHMWPE Sheets\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eHigh-temperature compression: Adequate amounts of nascent G1\u0026ndash;5 powders within a mold of 100 \u0026times; 100 \u0026times; (1\u0026ndash;2) mm were pre-compressed at 0.3 MPa and 128 \u0026deg;C for 400 s, pre-compressed at 2.5 MPa and 128 \u0026deg;C for 200 s, compressed at 10 MPa and 128 \u0026deg;C for 900 s, and then increased to 142\u0026ndash;145 \u0026deg;C for 400 s under a pressure of 0.3 MPa. The preformed samples were compressed at 2.5 MPa and 142\u0026ndash;145 \u0026deg;C for 200 s, the temperature was increased to 300 \u0026deg;C for 30 min under a pressure of 10 MPa, and the sample was compressed at 10 MPa and 300 \u0026deg;C for 2 h. Then, the preformed sheets in the molds were transformed into upper layers of a molding presser at a temperature of 124\u0026ndash;128 \u0026deg;C, isothermally crystallized for 3\u0026ndash;6 h under a pressure of 10 MPa, and slowly cooled to room temperature in the mold (Extended Data Fig. 10b).\u0026nbsp;\u003c/p\u003e\n\u003cp\u003ePreparation of sheets with 2D-ordered and 3D randomly constructed nanofibrils: Adequate amounts of nascent G1\u0026ndash;5 powders in a mold with a size of 100 \u0026times; 100 \u0026times; (1\u0026ndash;2) mm or 240 \u0026times; 240 \u0026times; (2\u0026ndash;40) mm were pre-compressed at 0.3 MPa and 128 \u0026deg;C for 200\u0026ndash;600 s, pre-compressed at 2.5 MPa and 128 \u0026deg;C for 400\u0026ndash;900 s, and compressed at 10 MPa and 128 \u0026deg;C for 300\u0026ndash;900 s. Then, the temperature was increased to 142\u0026ndash;145 \u0026deg;C or 145\u0026ndash;220 \u0026deg;C for 400 s\u0026ndash;1 h under a pressure of 0.3 MPa. The preformed samples were compressed at 2.5 MPa and 142\u0026ndash;145 \u0026deg;C or 145\u0026ndash;220 \u0026deg;C for 200\u0026ndash;900 s, and compressed at 10 MPa and 142\u0026ndash;145 \u0026deg;C or 145\u0026ndash;220 \u0026deg;C for 120 s\u0026ndash;2 h. Then, the temperature was decreased to 128 \u0026deg;C for 30 min under a pressure of 10 MPa, or the preformed sheets in molds were transformed into the upper layers of a molding presser at a temperature of 128 \u0026deg;C. The sheets were subsequently isothermally crystallized for 30 min\u0026ndash;23 h under a pressure of 10 MPa. For microscale 2D ordering at 145\u0026ndash;150 \u0026deg;C, the pre-formed sheets in molds were quenched in the upper layers at 20\u0026ndash;60 \u0026deg;C. For macroscale random 3D construction of nanofibrils, the compression molding temperatures were set to 145\u0026ndash;220 \u0026deg;C, and the preformed sheets in the molds were quenched in the upper layers at 20\u0026ndash;80 \u0026deg;C under a pressure of 10 MPa (Extended Data Fig. 10b).\u0026nbsp;\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eDynamic Nano-fibrillation of Nascent G1 Powders in POE Matrix\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003ePremixing and nano-fibrillation of G1 in POE1\u0026ndash;2: A batch internal mixer (HAAKE PolyLab OS Rheo-Drive 4, Thermo-Scientific Company) is used to premix or melt blend G1 with POE1\u0026minus;2. The rotation speed was 60/120 rpm, the mixing temperature was 80/115 \u0026deg;C, and the mixing time was 6 mm (Extended Data Fig. 10a).\u003c/p\u003e\n\u003cp\u003eCompression molding: Pellets or blends are pre-compressed in the mold at 0.3 MPa and 110/160 \u0026deg;C for 300 s, pre-compressed at 0.3 MPa and 110/160 \u0026deg;C for 100 s, pre-compressed at 2.5 MPa and 110/160 \u0026deg;C for 200 s, and compressed at 10 MPa and 110/160 \u0026deg;C for 300 s. The preformed samples are quenched in the upper layers of a molding presser at 25\u0026ndash;50 \u0026deg;C for 5 min (Extended Data Fig. 10b).\u003c/p\u003e\n\u003cp\u003eExtruded nano-fibrillation of G1 in POE1: A laboratory-scale corotating twin-screw extruder (HAAKE MiniLab Rheomex CTW5, Thermo-Scientific Company) is used to compound solid-state premixed POE1/G1 under cycling conditions for 5\u0026ndash;15 min. The weight composition ratio of G1 to POE is 95:5 wt%. The temperatures of the distinct zones are 115/135 \u0026deg;C. The rotation speed of the motor is 80\u0026ndash;300 rpm (Extended Data Fig. 10c).\u0026nbsp;\u003c/p\u003e\n\u003cp\u003eShear-induced orientation of G1 meso-particles into nanofibrils in POE1\u0026ndash;2: A laboratory-scale HAAKE MiniJet Pro combined with a corotating twin-screw extruder (HAAKE MiniLab Rheomex, Thermo-Scientific Company) is used to melt compounding and injection-molding of the pressed POE1-2/G1 nanocomposites. The weight composition ratios of G1 to POE1-2 are (100/99/97.5/95): (0/1/2.5/5) wt%. The extrusion/injection/molding temperatures are 160/190 \u0026deg;C and 40 \u0026deg;C. The rotation speed of the extruder was 80 rpm.\u0026nbsp;\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eReferences\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003e41. Aizezi, M. et al. High-strength triple shape memory elastomers from radiation-vulcanized polyolefin elastomer/polypropylene blends. \u003cem\u003eACS Appl. Polym. Mater.\u003c/em\u003e \u003cstrong\u003e1\u003c/strong\u003e, 1735\u0026ndash;1748 (2019).\u003c/p\u003e"}],"fulltextSource":"","fullText":"","funders":[],"hasAdminPriorityOnWorkflow":false,"hasManuscriptDocX":true,"hasOptedInToPreprint":true,"hasPassedJournalQc":"","hasAnyPriority":true,"hideJournal":true,"highlight":"","institution":"","isAcceptedByJournal":false,"isAuthorSuppliedPdf":false,"isDeskRejected":"","isHiddenFromSearch":false,"isInQc":false,"isInWorkflow":false,"isPdf":false,"isPdfUpToDate":true,"isWithdrawnOrRetracted":false,"journal":{"display":true,"email":"[email protected]","identity":"researchsquare","isNatureJournal":false,"hasQc":true,"allowDirectSubmit":true,"externalIdentity":"","sideBox":"","snPcode":"","submissionUrl":"/submission","title":"Research Square","twitterHandle":"researchsquare","acdcEnabled":true,"dfaEnabled":false,"editorialSystem":"","reportingPortfolio":"","inReviewEnabled":false,"inReviewRevisionsEnabled":true},"keywords":"","lastPublishedDoi":"10.21203/rs.3.rs-5814738/v1","lastPublishedDoiUrl":"https://doi.org/10.21203/rs.3.rs-5814738/v1","license":{"name":"CC BY 4.0","url":"https://creativecommons.org/licenses/by/4.0/"},"manuscriptAbstract":"During the past 70 years, it has been well established that melts of crystalline homopolymers adopt disordered random coil conformations and are densely packed in globules above the melting temperature (\u003ci\u003eT\u003c/i\u003e\u003csub\u003em\u003c/sub\u003e). Flow-induced melt orientation promotes crystal nucleation and polymorphism upon cooling. Here, we present polymerization-induced nano-to-mesoscale condensed-ordering as well as order-to-order and order-to-disorder (O-D) melt transformations at \u003ci\u003eT\u003c/i\u003e\u003csub\u003eO-D\u003c/sub\u003e in engineering polyethylene (E-PE). Moreover, we demonstrate fast chain explosion and re-entanglements at \u003ci\u003eT\u003c/i\u003e\u003csub\u003eO-D\u003c/sub\u003e. Liquid-crystalline-like macroscale 2D-ordered nanofibrils and meso-ordered superstructures self-assemble from high-order condensed states and ordered melts via isothermal crystallization, respectively. Self-reinforced optical and shape memory sheets are made by preserving nascent meso-ordering. Eplastomers, which exhibit both elastomeric and soft‒plastic properties, are prepared by nano-fibrillation of non-crosslinked E-PE meso-particles in a polyolefin elastomer matrix and dynamic exchange of 1D-ordered nanofibrillar networks with disordered 3D meso-granules. These findings enable further development of ultrahigh-performance intelligent, photonic and sustainable polymeric materials.","manuscriptTitle":"Polymerization-Induced Nano-Order, Ordered Melt, Chain Explosion and Meso-Superstructures of E-Polyethylene","msid":"","msnumber":"","nonDraftVersions":[{"code":1,"date":"2025-01-15 05:29:34","doi":"10.21203/rs.3.rs-5814738/v1","editorialEvents":[],"status":"published","journal":{"display":true,"email":"[email protected]","identity":"researchsquare","isNatureJournal":false,"hasQc":true,"allowDirectSubmit":true,"externalIdentity":"","sideBox":"","snPcode":"","submissionUrl":"/submission","title":"Research Square","twitterHandle":"researchsquare","acdcEnabled":true,"dfaEnabled":false,"editorialSystem":"","reportingPortfolio":"","inReviewEnabled":false,"inReviewRevisionsEnabled":true}}],"origin":"","ownerIdentity":"8cb87787-f796-458b-8800-2c68687301f5","owner":[],"postedDate":"January 15th, 2025","published":true,"recentEditorialEvents":[],"rejectedJournal":[],"revision":"","amendment":"","status":"posted","subjectAreas":[{"id":42839804,"name":"Physical sciences/Materials science/Nanoscale materials/Molecular self-assembly"},{"id":42839805,"name":"Physical sciences/Nanoscience and technology/Nanoscale materials/Structural properties"},{"id":42839806,"name":"Physical sciences/Chemistry/Polymer chemistry/Mechanical properties"},{"id":42839807,"name":"Physical sciences/Chemistry/Polymer chemistry/Nanocomposites"},{"id":42839808,"name":"Physical sciences/Physics/Atomic and molecular physics/Macromolecules and clusters"}],"tags":[],"updatedAt":"2025-01-16T12:40:16+00:00","versionOfRecord":[],"versionCreatedAt":"2025-01-15 05:29:34","video":"","vorDoi":"","vorDoiUrl":"","workflowStages":[]},"version":"v1","identity":"rs-5814738","journalConfig":"researchsquare"},"__N_SSP":true},"page":"/article/[identity]/[[...version]]","query":{"redirect":"/article/rs-5814738","identity":"rs-5814738","version":["v1"]},"buildId":"8U1c8b4HqxoKbykW_rLl7","isFallback":false,"isExperimentalCompile":false,"dynamicIds":[84888],"gssp":true,"scriptLoader":[]}